Effect of Preparation Parameter on Microstructure and Grain Refining Behavior of In Situ AlN-TiN-TiB₂/Al Composite Inoculants on Pure Aluminum

Abstract

The formation of in situ AlN-TiN-TiB₂/Al composite inoculants, which contain multi-phase refiner particles including AlN, TiN, TiB₂, Al₃Ti, and α-Al, was investigated using nitrogen gas injection. The effects of the main preparation parameters such as nitriding temperature, nitriding time, Ti content in melts, on the microstructure and grain refinement of in situ AlN-TiN-TiB₂/Al composite inoculants were studied. The shape, content and size of different ceramic particles in the inoculants can be tuned by controlling the nitriding temperature and time, inducing excellent refining and reinforcing effects on pure aluminum. As a result, the average grain size of pure aluminum can be reduced to about 122 ± 22 μm from original 1010 ± 80 μm by adding 0.3 wt % inoculants. The mechanical properties including the tensile strength, yield strength and microhardness of the refined as-cast pure aluminum are also improved.

1. Introduction

Inoculation treatment is a traditional casting technology to refine the microstructure of castings in the field of aluminum alloys. By adding some special master alloys (intermediate alloys) that contain a lot of particles serving as the substrates for heterogeneous nucleation, the grain size of the inoculated alloys can be obviously reduced and the comprehensive properties are improved [1,2,3,4]. In fact, most of the master alloys belong to Al-based composites, in which the aluminum-based solid solution is the matrix and the refiner particle is the secondary phase. Here, the secondary phase is generally a ceramic particle, which plays a role of grain refining. For example, in the conventional Al-5Ti-1B inoculants, TiB₂ and Al₃Ti particles provide the substrates for the nucleation of Al crystals. However, the activity of Al₃Ti is strongly affected by the chemical elements contained in an inoculated Al melt. When the Al melt contains special elements, such as Zr, V, and Mn, the performance of Al-5Ti-1B inoculants can sharply deteriorate, namely, with a poisoning effect of Al-5Ti-1B inoculants [5,6]. Moreover, when the concentration of silicon in Al-Si alloys exceeds 3 wt %, the refining effect of Al-5Ti-1B inoculants is greatly reduced due to the formation of Ti₅Si₃ that has a poor crystallographic matching with aluminum matrix [7]. In addition, while improving the amount of TiB₂ particles for more substrates, TiB₂ particles are inclined to agglomerate and precipitate in aluminum melt, which triggers the fading of refinement effect. To change the limitation of Al-5Ti-1B inoculants, other ceramic particles such as CeB₆ [8], LaB₆ [9], TiN [10,11], and B₄C [12], were also introduced to the study of inoculants. In particular, some multi-phase refiners show excellent refining effects on a variety of aluminum alloys, due to the presence of different refiner particles [13]. To sum up, for both single-phase refiners and multi-phase refiners, the amount and distribution of refiner particles play important roles for the inoculation behavior of composite inoculants [14]. Therefore, the preparation and processing of inoculants alloys, which determines both macro and micro structure of the inoculants, are very important and should be studied.

Recently, in situ AlN-TiN-TiB₂/Al composite inoculants were developed via the gas injection method in our previous work [15]. However, the influence of typical preparation parameters on the grain-refining effect of the inoculants was not described in detail. In this work, five in situ AlN-TiN-TiB₂/Al composite inoculants were prepared via gas injection method to research the influence of the nitriding temperature, nitriding time, Ti content and stirring device on the microstructure and grain refinement of in situ AlN-TiN-TiB₂/Al composite inoculants.

2. Experimental Procedures

Commercial Al-5Ti-1B alloy and titanium sponge with a purity of <99.5% were used as the starting materials. A WK-II-type non-consumable vacuum arc furnace (WuKe Optoelectronics Technology Ltd., Beijing, China) was employed to prepare the as-cast Al-8Ti-1B master alloy. Then, the as-received Al-5Ti-1B alloy and the as-cast Al-8Ti-1B alloy were remelted and treated by nitriding equipment, as shown in Figure 1. The experiment was carried out in a crucible-type resistance furnace (ZhongHuan experimental electric furnace Ltd., Tianjin, China). A highly pure graphite crucible with a graphite lid was used to melt Al-5Ti-1B alloys and Al-8Ti-1B alloys, respectively. The experimental process was as follows: Firstly, the alloy charges of approximately 100 g were heated in the crucible type resistance furnace. When the temperature of the melts reached 1173 K or 1373 K, highly pure nitrogen (purity: 99.997%, O₂ content <6 ppm) was introduced into the melts through the gas-guide tube (Al₂O₃ tube) for 1 h to 2 h at a gas flow rate of 1.0 L/ min. A stirring device was applied to mix the in situ synthetic ceramic particles with the Al melts, homogeneously. The detailed processing parameters for each composite inoculant are listed in

Table 1. Processing parameters for the inoculants.

Alloys No.	Composition	Nitridation Temperature (K)	Nitridation Time (h)	Whether the Stirring
A1	Al5Ti1B	1173	1	Stirring
A2	Al5Ti1B	1373	1	Stirring
A3	Al8Ti1B	1373	1	Stirring
A4	Al8Ti1B	1373	2	Stirring
A5	Al5Ti1B	1373	1	Without Stirring

. Finally, the melts after nitridation were then poured into a steel mold, which has a cylindrical cavity with a diameter of 20 mm and a length of 150 mm.

The specimens used for microstructure observation were cut from the surface and center of the ingot, respectively. The specimens were polished using standard metallographic techniques and etched with 0.5 vol % hydrofluoric acid after mechanical polishing. Microstructures and components of specimens were observed and analyzed by Philips XL 30 TMP scanning electron microscopy (SEM, Royal Dutch Philips Electronics Ltd., Amsterdam, The Netherlands) with energy-dispersive spectrometry (EDS, Royal Dutch Philips Electronics Ltd., Amsterdam, The Netherlands). Philips X′Pert MPD X-ray diffraction (XRD, Royal Dutch Philips Electronics Ltd, Amsterdam, The Netherlands) was used to identify the phase composition of specimens.

The in situ AlN-TiN-TiB₂/Al composites were used as inoculants and added into pure aluminum melts (purity: 99.75%) with an amount of 0.3 wt %. The melts were then stirred with a graphite rod for 30 s. After that, the melts were covered by the flux and hold for 10 min. Then the melts were poured into a steel mold with a size of 100 mm × 100 mm × 150 mm, which had a cylindrical cavity with a diameter of 20 mm and a length of 150 mm. Each metallographic specimen was cut from the middle part of the as-cast rod and polished to remove scratches. The specimens were etched by a reagent (75 mL HCl + 25 mL HNO₃ + 5 mL HF) for macrostructure observation. The grain size analysis was carried out using the linear intercept method. The tensile specimens were machined following standard ASTM E399 and the gauge length was about 50 mm with a diameter of 10 mm. A tensile test was performed on a SHT5305 type microcomputer-controlled electro-hydraulic servo universal testing machine (DunJie Measurement Devices Ltd, Wuhan, China) at a strain rate of ~1 × 10⁻³ s⁻¹. An HXD-1000 microhardness tester (Shimadzu Corporation, Kyoto, Japan) was used to determine the Vickers microhardness of specimens.

3. Results and Discussion

3.1. Effects of Nitriding Temperature, Nitriding Time, Composition, and Stirring Device on the Formation of In Situ AlN-TiN-TiB₂/Al Composites

The XRD analysis of five in situ AlN-TiN-TiB₂/Al composite inoculants are shown in Figure 2, and Figure 2b is the amplification of Figure 2a from 32° to 39°. It can be seen that sample A1 consists of α-Al, Al₃Ti, and TiB₂ phases, while sample A2, A3, A4, and A5 consist of α-Al, Al₃Ti, TiB₂, AlN, and TiN phases. However, the detailed fraction of Al₃Ti, AlN, and TiN phases in A2, A3, A4, and A5 is significantly different. In Figure 2a A1, there are α-Al, A1₃Ti and TiB₂ peaks in the XRD pattern of the Al-5Ti-1B melt treated at 1173 K for 1 h, whereas TiN and AIN peaks were not found. It is implied that the content of TiN and AIN phases is too little. With the increase of the treatment temperature from 1173 K to 1373 K, the relative intensity of TiN and AIN peaks increases slightly, whereas the A1₃Ti peaks are decreased as shown in Figure 2a A2. As Ti content increases from 5 wt % to 8 wt %, the intensity of A1₃Ti, TiN and AIN peaks significantly increases as shown in Figure 2a A3. In addition, as the treatment time is increased from 1 h to 2 h, the relative intensity of TiN and AIN peaks increase as shown in Figure 2a A4. To sum up, the relative intensity of TiN and AlN peaks increase with the increase of nitridating temperature, nitridating time and Ti content in melt.

If the stirring device is not used, the upper part of the Al-Ti-B melt forms a hard shell during nitridation at 1373 K, while the lower part of the melts remains good fluidity. The hard shell (A5) was peeled off for XRD analysis, as shown in Figure 2 A5. It can be seen that the hard shell consist of massive AlN phase.

The microstructures and EDS analysis of in situ AlN-TiN-TiB₂/Al composite inoculants are shown in Figure 3. From Figure 3a,b, it can be seen that sample A1 is composed of blocky-like Al₃Ti phase and a few TiB₂, AlN and TiN particles. The size of Al₃Ti phases is about ~50 μm and the size of TiB₂, AlN, and TiN particles is about ~1 μm. When the nitriding temperature is increased to 1373 K, the edge of blocky-like Al₃Ti phases occurs bursting crack. Meanwhile, the content of AlN and TiN particles also increases slightly as shown in Figure 3d,e. From Figure 3g,h, it can be seen that A3 composite, which was synthesized by Al-8Ti-1B alloy at 1373 K with the nitriding time of 1 h, is composed of massive blocky-like Al₃Ti phase and massive TiB₂, AlN and TiN particles. The edge of blocky-like Al₃Ti phases crack seriously in A3 composite and the size of Al₃Ti phases decreases. The large increase in the content of blocky-like Al₃Ti phase and TiB₂, AlN and TiN particles are all ascribed to the increase of Ti content in alloy. As it was known, the shape of TiN, AlN, and TiB₂ is square, hexagon and irregular, respectively [13]. Therefore, based on a combining consideration of the XRD results (Figure 2), the morphologies (Figure 3h,k) and the EDS results (Figure 3f,i,l), these particles can be confirmed as AlN, TiN, and TiB₂ phases. When nitriding time increases to 2 h, the size of Al₃Ti phase decrease to ~14 μm and a large number of TiN, AlN and TiB₂ particles are homogeneously formed in the matrix with a size of ~1 μm as shown in Figure 3j,k. Figure 3m,n shows the SEM images of A5, namely, the hard shell in the upper part of the crucible. Combined with the XRD (Figure 2e) and EDS analysis (Figure 3o), it can be seen that the appearance of AlN phase shows a whisker-like shape, very different from the other four. In particular, the AlN whiskers are agglomerated to AlN binds with sizes ranged from nano-scale to micron scale due to the enormous surface energy.

3.2. Thermodynamics Analysis and the Formation Mechanism

The following Reactions (1)–(4) would probably take place in Al-Ti-B melts during the process of nitrogen injection:

$$2\mathrm{Al}+{\mathrm{N}}_2\to 2\mathrm{AlN}$$

(1)

$$[\mathrm{Ti}]+\frac{1}{2}{\mathrm{N}}_2\to \mathrm{TiN}$$

(2)

$${\mathrm{Al}}_3\mathrm{Ti}\to 3[\mathrm{Al}]+[\mathrm{Ti}]$$

(3)

$$3[\mathrm{Al}]+\frac{3}{2}{\mathrm{N}}_2\to 3\mathrm{AlN}$$

(4)

$$2\mathrm{Ti}+{\mathrm{N}}_2\to 2\mathrm{TiN}$$

(5)

where Al and Ti are solid Al and Ti, [Al] and [Ti] are free atoms in the melts (or liquid Al and Ti), N₂ is nitrogen gas, and AlN, TiN, and Al₃Ti are all in solid state.

The temperature dependence of Gibbs free energy changes for Reactions (1),(2),(5) are shown in Figure 4. The lines for Reactions (1) and (5) were directly obtained from the formulas in [16]. From the Gibbs free energy changes, the possibility to form TiN is higher than that of AlN because the Gibbs free energy values of Reaction (5) (TiN) is smaller than Reaction (1) (AlN) as shown in Figure 4. However, the state of Ti in aluminum melt might be none-solid Ti. Thus, there are many titanium atoms ([Ti]) in aluminum melts, which mainly depends on the amount of Ti and the temperature. It is very different from the standard state. The standard Gibbs free energy changes for Ti from puresolid to 1 wt % solute in aluminum melts is [17]:

$$\mathrm{Ti}\left(\mathrm{s}\right)\to [\mathrm{Ti}] \Delta G=-97116-13.624T$$

(6)

When combining Reactions (2) and (6), the Gibbs free energy change of the reaction of [Ti] with N₂, which is as the standard state of 1 wt % Ti in aluminum, can be obtained as shown the line for R2 in Figure 4. It can be seen that the possibility of the formation of TiN by Reaction (2) is lower than that of AlN by Reaction (1) as shown in Figure 4. This result is in accord with the above experimental result, which is suggested that AlN fraction is much higher than TiN.

According to the binary Ti-Al alloys phase diagram as shown in Figure 5, Al-5Ti-1B melts at 1173 K and Al-8Ti-1B melts at 1373 K are located in the two-phase region (L + TiAl₃), so there is equilibrium between phases, which can be expressed as:

$$3{[\mathrm{Al}]}_{\mathrm{L}}+{[\mathrm{Ti}]}_{\mathrm{L}}\rightleftharpoons {\mathrm{Al}}_3{\mathrm{Ti}}_{\mathrm{S}}$$

(7)

When [Ti] react with N₂ to form TiN, the concentration of [Ti] in the nitriding Al melt significantly decreases. Then, the reduction of [Ti] would destroy above phase equilibrium (7) and make it develop toward the reverse reaction of (7), leading to the reduction of Al₃Ti phase.

At high temperature, the amount of solid Al₃Ti is quite small and [Ti] is mainly dissolved in the Al melt. As the melts are cooled, Al₃Ti phase is precipitated from the melt and grows into blocky particle. When N₂ is introduced to the melt with high Ti content, [Ti] is mainly expended by nitridation reaction such that Al₃Ti phase cannot grow into a large strip to make the ceramic particles aggregate.

Based on the above results, Figure 6 shows the schematic of the reactions between N₂ and Al-Ti-B melts. AlN and TiN particles can be directly formed from the reactions during gas injection. The volume fraction of AlN and TiN particles is small during nitrogen injection to Al-5Ti-1B melts at 1173 K for 1 h because of low nitriding temperature, short nitriding time, and low Ti content (as shown in Figure 6a). The volume fraction of AlN and TiN particles dramatically increases during nitrogen injection to Al-8Ti-1B melts at 1373 K for 2 h (as shown in Figure 6b). The volume fraction of AlN and TiN particles increases with the increase of nitriding temperature, nitriding time, and the amount of [Ti] in Al-Ti-B melts. The nitriding reaction mainly takes place at the interface of gas/melt, and the AlN and TiN particles move upwards together with the gas bubbles. The product particles are detached from the interface of gas/melts and enter into the melts due to continuous stirring. When the stirring device is not used, the AlN and TiN particles move upwards together with the gas bubbles and the moving gas bubbles can also absorb TiB₂, AlN, and TiN particles in the melts and move upward together with them. Therefore, the particles can grow up with the continuous deposition of AlN or TiN as shown in Figure 6c.

3.3. Refining and Reinforcing Effects of the Inoculants on Pure Aluminum

Figure 7 shows the macrostructure of the as-cast pure aluminum sample that is refined by different inoculants with the same addition amount 0.3 wt %. The pure aluminum sample without inoculation exhibits coarse columnar grains. The five in situ AlN-TiN-TiB₂/Al composites all show obvious grain refining effects on pure aluminum, whereas the refining efficiencies are not quite similar. The detailed grain size analysis is shown in Figure 8. The mean grain size of as-cast pure aluminum samples without inoculation is about 1010 ± 80 μm. Based on the quantitative grain size analysis (Figure 8), the grain size of the sample becomes smaller and smaller when the 0.3 wt % A1, 0.3 wt % A2, 0.3 wt % A3, and 0.3 wt % A4 are added into pure aluminum in turn. Interestingly, by adding 0.3 wt % A4, the big grain is obviously refined into small grains with an average of 122 ± 22 μm. It is suggested that the refining effect of in situ AlN-TiN-TiB₂/Al composite inoculants increases with the increase of the volume fraction of AlN and TiN particles. Moreover, such small adding content with significant refining effect suggests A4 as economic inoculants for aluminum alloys.

The excellent grain refining effect of in situ AlN-TiN-TiB₂/Al composites can be attributed to the following two factors: At the first, Al₃Ti phases become smaller, which provides more substrates for nucleation of α-Al due to the increased specific areas [18,19]. Secondly, a large number of AlN and TiN particles are generated during nitriding besides Al₃Ti phase and TiB₂ particles and these particles dispersed homogeneously in the matrix. This is beneficial for refining aluminum grains. The α-Al phase and TiN phase belong to the face-centered cubic structure. The mismatch degree of the close-packed plane can be expressed by the following equation:

$$\mathsf{\delta }=\frac{\left|d_m-d_p\right|}{d_p}$$

(8)

where d_m is the close-packed spacing of matrix phase, d_p is the close-packed spacing of the precipitate phase. According to the planar lattice mismatch theory presented by Bramfitt [20], the nucleation particles are the most effective with the mismatch of the two phases is less than 6%. The mismatch degree (δ) of several possible coherent interfaces between TiN and Al are shown in

Table 2. Mismatch degree of possible coherent interfaces between TiN and Al.

Number	TiN		Al		δ/%
	d/nm	(hkl)	d/nm	(hkl)
1	0.2440	111	0.2330	111	4.72
2	0.2120	200	0.2020	200	4.95
3	0.1496	220	0.1430	220	4.61
4	0.1227	311	0.1220	311	0.57
5	0.1223	320	0.1170	222	4.53

For AlN phase, the crystal structure is close-packed hexagonal (HCP), different from that of Al.

. Based on the calculation results, the mismatch degree between α-Al phase and TiN phase is less than 6%, so it can be used as the most effective nucleating nuclei for the aluminum alloy.

The lattice constants of Al and AlN are a_Al (4.049 × 10⁻¹⁰ m) and a_AlN (3.113 × 10⁻¹⁰ m), c_AlN (4.981 × 10⁻¹⁰ m), respectively. It can be known that $a_H/a_F$ = 0.76876, $c_H/a_H$ = 1.6. According to the edge-to-edge matching crystallographic model [21], the match of crystal structures between FCC/HCP systems can be used to calculate the misfit of AlN/Al [22]. In this model, it is reasonable to use 10% and 6% as a critical value for the interatomic spacing misfit and the interplanar spacing misfit. Here, the misfit of ${[11\overline{2}0]}_H$//${[110]}_F$, ${\left(1\overline{1}01\right)}_H$//${\left(1\overline{1}1\right)}_F$ is 1.42% < 6%, 8.73% < 10%, respectively. They both meet the requirements of the critical value. Thus, it is suggested that in situ AlN particles are potential nucleation substrates of aluminum. Furthermore, Cui et al. [23] also found that there existed a crystal parallel relationship between AlN/Al: ${\left(1\overline{1}01\right)}_{\mathrm{AlN}}$//${\left(1\overline{1}1\right)}_{\mathrm{Al}}$. Figure 9 shows the crystal structures of Al, TiN and AlN. Figure 9b is the crystallographic model of Al and AlN. Thus, it is confirmed that ${[110]}_{\mathrm{Al}}$ in ${\left(1\overline{1}1\right)}_{\mathrm{Al}}$ can nucleate and grow up along ${[11\overline{2}0]}_{\mathrm{AlN}}$ in ${\left(1\overline{1}01\right)}_{\mathrm{AlN}}$.

Figure 10 shows interface microstructure of in situ formed TiN, AlN, and TiB₂ particles. From Figure 10a it can be found that the ceramic particles have three different size of micron-scale, submicron-scale, and nanometer-scale. Figure 10b shows the interface feature between hexagonal AlN phase and α-Al phase. Figure 10c shows the interface feature between α-Al phase and in situ formed TiN phase, in which a clean interface is observed. Thus, a strong interfacial bonding of the matrix and TiN particles can be obtained. Figure 10d shows the interface feature between α-Al phase and TiB₂ phase. It is found that an Al₃Ti layer is formed on the surface of TiB₂ particles during the grain refiner production process. This phenomena may significantly improve the potency of TiB₂ for nucleation of the α-Al, according to the reported literature [24].

However, the grain size of the sample becomes big when 0.3 wt % A5 is added into pure aluminum. The refining effect of A5 is not obvious because A5 contains a large volume fraction of AlN whiskers in addition to small amounts of Al₃Ti, TiB₂, and TiN, suggesting that AlN whiskers have no effect in the process of refining. That is because AlN whiskers have a light specific weight, larger specific surface, and strong adsorption, so they are likely to be floating on the surface of the Al melt.

Figure 11 shows the microhardness values of the as-cast Al samples after inoculation with different inoculants. The microhardness values becomes higher and higher with the increased sample number from A1–A4. However, the microhardness values decreases when 0.3 wt % A5 is added into pure aluminum. In particular, the mcirohardness of alloys refined by adding 0.3 wt % A4, has reached up to 107.88 HV, which is improved 162% than that of pure Al without inoculation.

Figure 12 shows the mechanical properties of as-cast pure aluminum samples inoculated with different inoculants. The tensile strength, yield strength, and elongation of pure Al are 41.9 MPa, 28.5 MPa and 37.2%, respectively. The tensile strength and the yield strength of the specimens are both increased, but the elongation is decreased with A1, A2, A3, and A4 inoculant additions. It can be seen that the tensile strength and yield strength increase rapidly up to 109.7 MPa and 77.2 MPa with 0.3 wt % A4 inoculant. Although the elongation of the inoculated materials has a certain decline, the finally value is still in an acceptable level. However, the tensile strength and yield strength decreases as 0.3 wt % A5 is added into pure aluminum. It is suggested that the addition of AlN whiskers do not show an obvious reinforcing effect on pure aluminum. In our opinion, AlN whiskers are light-weight and have larger specific surface and strong adsorption effects, resulting in the whiskers floating on the surface of Al melt. Therefore, AlN whiskers are difficult to be added in the Al melt.

The improved mechanical properties of pure aluminum modified by in situ AlN-TiN-TiB₂/Al composite inoculants are mainly attributed to the grain refining and particle reinforcing effects. On one hand, according to the Hall-Petch formulation, the yield strength of the composite materials can be expressed as the following equation:

$${\mathsf{\sigma }}_y={\mathsf{\sigma }}_i+k_y{d}^{-\frac{1}{2}}$$

(9)

where σ_y is the yield strength, σ_i is a frictional stress resisting the motion of gliding dislocations or an internal back stress, k_y is the Hall-Petch slope which reflects resistance of grain boundary to slip transfer and d is the average grain diameter. From R9, the yield strength of the material is related to the grain size. The smaller the grain size reduced, the higher yield strength can be obtained. In addition, the finer the grain size, the more number of the grain boundaries exist. The grain boundaries can impede dislocation movement and result in the significant improvement of strength of the composite materials, which is called grain boundary strengthening [25]. On the other hand, numerous fine ceramic particles, such as TiB₂, TiN, and AlN particles, are uniformly distributed in the aluminum matrix and, thus, play an important role of dispersion strengthening. Fine ceramic particles can pin the dislocations and hinder dislocation movement to improve the deformation resistance stress which can ultimately lead to the enhancement of the strength. However, the finer grain size not only makes the material have high hardness and strength [26], but also gives it good plasticity, which means good comprehensive mechanical properties. However, particle reinforcement still leads to a decrease of the elongation of the composites.

4. Conclusions

(1)

A multiple phase inoculant, the in situ AlN-TiN-TiB₂/Al composite containing AlN, TiN, TiB₂, Al₃Ti and α-Al, was successfully prepared by nitrogen gas injection. The volume fraction, size, and distribution of TiN and AlN particles can be controlled by tuning the preparation parameters, such as nitriding temperature, nitriding time, Ti content in melt, and stirring.

(2)

The refining effect of in situ AlN-TiN-TiB₂/Al composite inoculants increases with the increase of the content of TiB₂, AlN, and TiN particles and the decrease of the size of the Al₃Ti phase. The average grain size of the pure aluminum can be reduced to 122 ± 22 μm from 1010 ± 80 μm by adding of 0.3 wt % A4.

(3)

The excellent grain refining effect of in situ AlN-TiN-TiB₂/Al composites is attributed to the miniaturization of Al₃Ti phase, the diversification and homogeneous distribution of ceramic particles, such as AlN, TiN, and TiB₂ particles caused by nitridation and stirring.

(4)

The mechanical properties of the pure aluminum have been obviously improved by the addition of in situ AlN-TiN-TiB₂/Al composite inoculants. The tensile strength, yield strength, and microhardness of the pure aluminum are increased from 41.9 MPa, 28.5 MPa, and 41.18 HV to 109.7 MPa, 77.2 MPa, and 107.88 HV.