Understanding the High-Temperature Deformation Behaviors in Additively Manufactured Al6061+TiC Composites via In Situ Neutron Diffraction

Abstract

Aluminum matrix composites (AMCs) are designed to enhance the performance of conventional aluminum alloys for engineering applications at both room and elevated temperatures. However, the dynamic phase-specific deformation behavior and load-sharing mechanisms of AMCs at elevated temperatures have not been extensively studied and remain unclear. Here, in situ neutron diffraction experiments are employed to reveal the phase-specific structure evolution of additively manufactured Al6061+TiC composites under compressive loading at 250 °C. It is found that the addition of a small amount of nano-size TiC significantly alters the deformation behavior and increases the strength at 250 °C in comparison to the as-printed Al6061. Unlike the two-stage behavior observed in Al6061, the Al6061+TiC composites exhibit three stages during compression triggered by changes in the interphase stress states. Further analysis of Bragg peak intensity and broadening reveals that the presence of TiC alters the dislocation activity during deformation at 250 °C by influencing dislocation slip planes and promoting dislocation accumulation. These findings provide direct experimental observations of the phase-specific dynamic process in AMCs under deformation at an elevated temperature. The revealed mechanisms provide insights for the future design and optimization of high-performance AMCs.

1. Introduction

Aluminum alloys have long been essential for the automotive, aerospace, marine, and other industries due to their high strength-to-weight ratio, low cost, and excellent corrosion resistance [1,2,3]. However, common aluminum alloys often exhibit degraded mechanical performance at elevated temperatures due to their low melting points, thermal softening, microstructural instability, and low creep resistance [4,5,6]. Particle-reinforced aluminum matrix composites (AMCs) [7,8,9,10], which combine aluminum alloys as the matrix and secondary ceramic particles as the reinforcement phase, are designed to augment the performance capabilities of conventional Al alloys, particularly in environments where traditional aluminum alloys fall short, such as those at elevated temperatures. Unlike common high-strength Al alloys (e.g., Al2000, Al6000, and Al7000 series), which rely on solid-state metastable precipitation to achieve high strength [11,12,13,14], AMCs use thermodynamically stable ceramic particles as reinforcement, which enable robust strengthening at both room and elevated temperatures [15,16,17,18].

Numerous studies have demonstrated the strengthening effects in AMCs at both room and elevated temperatures by adding ceramic particles, such as TiC [19,20], SiC [21,22], Al₂O₃ [23,24], and others [25,26]. Adding particles activates several strengthening mechanisms to enhance the mechanical strengths of AMCs, including Orowan strengthening, the Hall–Petch effect, load sharing, etc. [27,28,29]. Load sharing between phases can exhibit complex multistage behavior throughout the deformation process, as observed in other dual or multiphase metal alloys [30,31]. In addition to load sharing between phases, the deformation behavior of the Al matrix differs significantly at elevated temperatures in terms of dislocation annihilation, the activation of additional dislocation slip systems, and precipitation-induced chemistry changes [32,33]. Adding particles can impede the dislocation motion of the Al matrix, potentially leading to more complex and distinctly different deformation behaviors in AMCs at elevated temperatures.

However, most previous research has focused on the collective response of bulk AMCs [10,34,35,36]. The dynamic phase-specific response at the grain level, especially during elevated-temperature loading, has not been well studied. This gap hinders a deeper understanding of deformation and load-sharing mechanisms as well as the design of AMCs with enhanced mechanical performance at elevated temperatures. Investigating the load-sharing and deformation behavior of AMCs at elevated temperatures requires a highly complex in situ setup, including mechanical loading, sample heating, and neutron or synchrotron diffraction. However, most previous studies on metal matrix composites using neutron or synchrotron X-rays have focused on room-temperature deformation [37]. The load-sharing and deformation behaviors of AMCs at elevated temperatures have not been extensively studied due to the limited availability of high-flux and high-resolution neutron and synchrotron sources [38,39,40].

In fusion-based additive manufacturing (AM) processes [41,42,43,44], the unique rapid heating and cooling cycles with localized high melt pool temperatures and rapid solidification rates [45,46,47,48,49] may significantly influence interphase bonding, dislocation density, and solidification texture. These factors potentially affect the load sharing, deformation behaviors, and mechanical properties of AMCs. In this work, additively manufactured Al6061+TiC composites were fabricated by the laser powder bed fusion (LPBF) process and used as a model material for in situ neutron diffraction studies to understand the phase-specific deformation behaviors at an elevated temperature. Neutrons’ capability of interacting with lightweight elements such as carbon as well as their deep-penetrating capability to measure the bulk sample on average give rise to a high-quality signal of the Al matrix, a small amount of TiC, and minor precipitate phases from the alloy, clearly characterizing the dynamic and kinetic evolution of phase stability, lattice strain, and phase-specific stress under combined sample environments, such as thermomechanical testing at elevated temperatures.

2. Materials and Methods

2.1. Materials

Al6061, Al6061+2 vol.%TiC, and Al6061+5 vol.%TiC were fabricated by the laser powder bed fusion (LPBF) AM process. The Al6061 powders used in the LPBF process were purchased from Valimet (Stockton, CA, USA). The detailed chemical composition of the Al6061 powders is shown in

Table 1. Chemical composition of as-received Al6061 powders.

Element	Al	Mg	Si	Cu	Cr	Fe	Ti	Mn	Zn	Other
Weight percentage	Balance	0.86	0.55	0.27	0.1	0.09	0.01	<0.01	<0.01	<0.15

. The Al6061+2 vol.%TiC and Al6061+5 vol.%TiC powders were prepared by the planetary ball milling of Al6061 powders and TiC nanoparticles (83 nm, SSNano, Houston, TX, USA). Alumina grinding jars with a capacity of 100 mL and alumina grinding balls with a diameter of 6 mm were used for this study. The jars were loaded with 20 g of Al6061+TiC powder, maintaining a ball-to-powder weight ratio of 5:1. The disc rotation speed was set to 200 RPM, with a 10 min pause after every 20 min milling cycle. The total milling time was 8 h.

The LPBF process employed a pulsed laser beam with a Gaussian beam profile. The detailed laser processing parameters are listed in

Table 2. Processing parameters for LPBF process.

Laser Power	Beam Diameter	Exposure Time	Point Distance	Hatching Spacing
200 W	70 µm	200 µs	80 μm	80 μm

. Printing was conducted using the stripe scan pattern [50], with the scanning vector rotated by 67° following each layer. The printed sample had dimensions of 15 mm (length), 7 mm (width), and 20 mm (height).

2.2. In Situ Neutron Diffraction

In situ neutron diffraction experiments were conducted at the VULCAN diffractometer, Spallation Neutron Source, Oak Ridge National Laboratory [51], to investigate the load-sharing and deformation behavior of Al6061 and Al6061+TiC AMCs at the elevated temperature of 250 °C. Figure 1 shows the experiment setup, and more details on instrument layout and optics can be found in reference [51]. The cuboid specimen, measuring 12 mm (length) × 6 mm (width) × 6 mm (height), was initially heated to 250 °C using induction coils and maintained at this temperature for over ten minutes to allow the temperature to stabilize. Subsequently, uniaxial compression tests were performed along the sample’s length using the MTS load frame until approximately 10% strain was reached or strain hardening ceased. At the same time, time-of-flight neutron diffraction was continuously conducted on a gauge volume of 5 mm × 5 mm × 6 mm. The neutron gauge volume was defined by neutron incident slits (with a 5 mm horizontal opening and a 6 mm vertical opening) and 5 mm receiving collimators (Figure 1b). The incident neutron beam was directed at a 45° angle relative to the loading direction. The diffracted neutron beams at −90° (Bank 1) and +90° (Bank 2) relative to the incident beam were captured by two neutron detectors to obtain diffraction data from the (hkl) crystal planes along the loading direction (LD) and the normal direction (ND). The sample building direction in the LPBF was aligned with the ND in the in situ neutron diffraction experiment. A K-type thermocouple was spot-welded to the bottom center of the sample (Figure 1a), which was used to monitor and control sample temperature via the PID controller. To investigate the effect of test temperature on deformation behavior, three additional Al6061 and Al6061+TiC samples were tested at room temperature for comparison with the results obtained at 250 °C.

2.3. In Situ Neutron Data Analysis

The in situ neutron data collected during the continuous loading and unloading were chopped into 5 min time bins using VDRIVE software (https://web.ornl.gov/~kean/VDRIVE/VDRIVE.html, accessed on 1 July 2024) [52]. Single-peak fitting (SPF) for the hkl peaks of both the Al and TiC phases was performed for each data bin using VDRIVE software to obtain hkl-specific diffraction peak properties. Rietveld refinement of the full diffraction patterns was conducted for each data bin using GSAS with EXPGUI (https://subversion.xray.aps.anl.gov/trac/EXPGUI, accessed on 1 July 2024) [53,54] to obtain the phase-averaged lattice information, the secondary Mg₂Si phase fraction, and the site occupancies of TiC.

The lattice strain evolution was calculated using two methods: (1) a method based on the hkl-specific d-spacing determined from the SPF, and (2) a method based on the average lattice parameter of each phase determined from the Rietveld refinement:

$$\epsilon ={\displaystyle \frac{d-d_0}{d_0}}$$

(1)

where $d$ is the d-spacing of hkl lattice planes or the lattice parameter of the Al or TiC phases during loading and $d_0$ is the reference d-spacing or lattice parameter before loading. The effects of thermal residual stress on d₀ determination were not considered, as previous studies have shown that no thermal residual stress exists between the Al and reinforcement phases at temperatures of 200 °C or higher [55].

On the basis of the lattice strain and Hooke’s law, the phase-specific stress was calculated by [56]:

$${\sigma }_{p,i,11}={\displaystyle \frac{E_i}{\left(1+{\nu }_i\right)\left(1-2{\nu }_i\right)}}\left[\left(1-{\nu }_i\right){\epsilon }_{i,11}+{\nu }_i ({\epsilon }_{i,22}+{\epsilon }_{i,33})\right]$$

(2)

where i represents TiC or Al, ${\sigma }_{p,i,11}$ is the phase-specific stress along LD, $E_i$ is the diffraction elastic constant (i.e., the elastic constant determined by the slope of the macroscopic stress versus the elastic lattice strain curve in the elastic stage during in situ neutron diffraction), and ${\nu }_i$ is Poisson’s ratio. Specifically, $E_{Al}$= 68 GPa, $E_{TiC}$= 203 GPa, and ${\nu }_{Al}$ = 0.31, ${\nu }_{TiC}$= 0.22. ${\epsilon }_{i,11}$ is the lattice strain along LD, ${\epsilon }_{i,22}$ and ${\epsilon }_{i,33}$ are the lattice strains along the ND and the transverse direction (TD). The ND and TD are two axes perpendicular to LD in the 3D Cartesian coordinate system. The lattice strain along the TD is assumed to be the same as that along the ND.

The hydrostatic stress was calculated by [57]:

$${\sigma }_{hy,i}={\displaystyle \frac{{\sigma }_{p,i,11}+{\sigma }_{p,i,22}+{\sigma }_{p,i,33}}{3}}$$

(3)

where ${\sigma }_{p,i,11}$, ${\sigma }_{p,i,22}$, and ${\sigma }_{p,i,33}$ are the phase-specific stress along the LD, ND, and TD, respectively. The deviatoric stress was calculated by [57]:

$${\sigma }_{d,i,11}={\sigma }_{p,i,11}-{\sigma }_{hy,i}$$

(4)

where ${\sigma }_{p,i,11}$ is the phase-specific stress and ${\sigma }_{hy,i}$ is the hydrostatic stress.

3. Results and Discussion

3.1. Engineering Stress–Strain

Figure 2 presents the compressive engineering stress–strain curves for the Al6061 and Al6061+TiC composites at 250 °C. The addition of TiC significantly enhances the yield strength, which increases from 66 MPa for Al6061 to 96 MPa for Al6061+2%TiC and further to 104 MPa for Al6061+5%TiC. Moreover, the hardening behavior during the plastic deformation is also altered with TiC presence. After unloading and reloading, the Al6061 sample exhibits significant softening, whereas the Al6061+TiC samples demonstrate continuous hardening at larger strains. The macroscopic curves indicate different microscopic mechanisms with and without TiC, which is explained later via analysis of in situ neutron diffraction. Although unloading and reloading were conducted in the experiment, the elastic behavior in these steps is not discussed in the following analysis of lattice strains and phase stresses, and the corresponding data points are not plotted in the figures for simplicity.

3.2. Phase Stability and Lattice Strains

With various solid solutes in the Al matrix as well as dispersive ceramics, the high temperature possibly destabilizes the solid solution and drives the atom diffusion between phases during mechanical testing. This chemistry change may lead to a different $d_0$ value in Equation (1), and it can thus influence the accurate determination of elastic lattice strain and phase-specific stress via the in situ diffraction approach [56]. It is thus essential to understand the phase stability first. Neutron diffraction probes the matrix Al phase, TiC ceramic, and minor precipitates, such as Mg₂Si [13] (Figure 3a), which benefit from its unique capabilities of deep penetration in a bulk-scale volume and the high sensitivity of lightweight elements. The evolution of the diffraction patterns during compression at 250 °C is visualized in Figure 3b, which reflects the phase-specific deformation behaviors and the phase stability.

Rietveld refinement was conducted to quantify the weight fraction evolution of the TiC and Mg₂Si in Al6061+5%TiC during deformation at 250 °C. Both the peak height in the contour plot of neutron diffraction (Figure 4a) and the weight fraction via refinement (Figure 4b) reflect a stable TiC weight fraction of ~8%. This value agrees with the nominal TiC composition (5% by volume or 8.7% by weight). The slight decrease in the TiC weight fraction after printing may be due to some TiC particles attaching to the ball milling media and thus being slightly lost during the ball milling process. By contrast, the weight fraction of Mg₂Si initially increases from 0.4% to 0.9% and then stabilizes (Figure 4b). This suggests that Mg₂Si does not fully precipitate during the LPBF process and continues to precipitate during the compressive test at 250 °C.

Rietveld refinement was further used to study the possible changes in the chemical composition of the TiC phase. During the LPBF process, the atom’s interdiffusion between phases may be activated. This could result in some Al or other solutes dissolving into TiC by partially replacing Ti in the as-printed AMCs. Here, the model formula (Ti_1−xAl_x)C is employed in the Rietveld refinement by giving freedom of composition variation at the cation site. On the basis of this, the “occupancy” of the Al at the Ti-site during the compressive test of Al6061+5%TiC at 250 °C was calculated to indicate the chemical stability of the TiC phase. The results reveal a small “occupancy” of Al in TiC, and the value did not significantly change at 250 °C (Figure 4c). Thus, the occupancy refinement again supports the conclusion of the chemically stable TiC ceramic phase at this elevated temperature in addition to the support from the unchanged weight fraction.

Therefore, the Al lattice strain during compression at 250 °C may result from both the mechanical deformation and the time-dependent chemical changes, while the TiC lattice strain solely reflects the mechanical response. In the analysis of phase stress in the following, the deviatoric components of stress are used to deconvolute the effects of chemistry variation in the phase [56]. Although precipitates like Mg₂Si are probed in neutron diffraction, the trivial fraction (<1%) is not considered to significantly impact the composite behavior over the TiC strengthening. The following analysis focuses on the Al matrix and the TiC phase.

To uncover the phase-specific mechanical responses, lattice strains for different lattice planes in the Al and TiC phases during the compressive loading of Al6061 and Al6061+TiC composites at 250 °C were determined through single-peak fitting of neutron diffraction patterns. In Al6061, the lattice strains across all Al lattice planes show an approximately linear increase throughout the loading process, encompassing both the elastic and plastic stages (Figure 5a). The divergence in lattice strains between different lattice planes is minimal, indicating low elastic and plastic anisotropy at 250 °C.

In Al6061+TiC composites, both the Al and TiC phases initially exhibit an approximately linear increase in lattice strain with applied stress (Figure 5b,c). The rate of increase is higher for the Al phase than for TiC, reflecting a higher elastic modulus of TiC than of Al. At the applied stresses of approximately −108 MPa for Al6061+2%TiC and −123 MPa for Al6061+5%TiC, the lattice strain for Al abruptly decreases, while the lattice strain for TiC continues to increase. This sudden change can be attributed to the yielding of the Al matrix, which transfers more load from the softer Al matrix to the harder TiC particles. The minimal divergence in lattice strain between different lattice planes for both Al and TiC phases indicates low elastic and plastic anisotropy in both phases. Given this low anisotropy, the average phase-specific lattice strains were determined from the changes in the average lattice parameter of each phase, which were calculated via Rietveld refinement of the full diffraction patterns. The comparisons between the average lattice strains and the {311} lattice strains via SPF are shown in Figure 5d–f and are consistent with each other. These average lattice strains were used for further phase stress analysis.

3.3. Phase-Specific Stress

The Al and TiC phase-specific stresses were directly calculated on the basis of the average lattice strain obtained from Rietveld refinement and Hooke’s law (Equation (2)) without consideration of the chemical changes. These “phase stresses” are plotted in Figure 6. In Al6061, the “stress” of the Al phase increases nearly linearly along the LD while remaining at a low minimal value along the ND (Figure 6a). This result aligns with the expected triaxial stress calculation during the uniaxial loading of a single-phase component.

In the Al6061+TiC composites (Figure 6b,c), the “stress” of both the Al and TiC phases initially shows a continuous increase with applied stress along the LD. Following the yielding of the Al matrix, the rate of “stress” increase in the Al phase decelerates in Al6061+2%TiC and even decreases in Al6061+5%TiC. In the TiC phase, the “stress” exhibits a higher rate of increase after the Al matrix yields. In addition to the overall trend, fluctuations were observed in “phase stress”. As discussed above, this may result from both the mechanical response and the chemical changes at the elevated temperature simultaneously. To deconvolute these two types of responses, the deviatoric and hydrostatic stress components of the phase-specific stress were calculated. The deviatoric component suppressed the effects of $d_0$ variation due to chemistry, temperature, and so forth, and this reflects the pure mechanical response in this complex environment; the hydrostatic component may reflect the comprehensive effects of all these factors [56].

Deviatoric stress is responsible for shearing and distortion in a material, which is directly related to yielding and plastic deformation. Adding TiC decreases the deviatoric stress of the Al phase. At an applied stress of −100 MPa, the deviatoric stress values of the Al phase in Al6061, Al6061+2%TiC, and Al6061+5%TiC are −73 MPa, −66 MPa, and −63 MPa, respectively. The reduction in deviatoric stress due to the presence of TiC particles delays the yielding of the Al6061+TiC composite, which explains the strength enhancement in the macro strain–stress curve in Figure 2.

Compared with the “phase stress”, the deviatoric stress shows a very smooth change in response upon applied stress (Figure 6d–f). This enables the detection of slight but important details. The Al6061+5%TiC sample that has a better signal of the TiC phase clearly exhibits three stages during the deformation, as distinguished by the slopes of the Al and TiC phases (Figure 6f). At the beginning of loading, both phases deformed elastically, showing the same deviatoric stress. After about −75 MPa of applied stress, another distinct stage is observed before the obvious plastic deformation in which the slopes differ between the phases. The TiC phase shows higher deviatoric stress than the Al phase. This may be attributed to the initiation of localized plastic deformation of the Al phase near TiC particles, leading to the early yielding of the material. The curve divergence increases upon loading, and it is further amplified in the third stage, in which the plastic deformation occurs over the Al phase globally. It is clearly observed that the existence of TiC shares the deviatoric stress and significantly reduces the deviatoric stress in the Al matrix in the plastic region under high applied stress. This can greatly delay the failure occurrence in the Al phase. In the Al6061+2%TiC sample, the slopes of the deviatoric stresses exhibit trends consistent with those in the Al6061+5%TiC sample, although the data of TiC have larger fitting error bars (Figure 6e).

The “hydrostatic stress” of the Al phase initially starts at zero and becomes increasingly negative in all samples (Figure 6g–i), indicating a continuous increase in compressive hydrostatic stress during loading. Compared to the deviatoric stress, the “hydrostatic stress” of Al may result from both the average hydrostatic stress state throughout the Al–TiC phase boundary to the depth of the Al grains and the change in chemical composition due to precipitate formation (Figure 4a,b). Differently, as the TiC phase is chemically stable, the “hydrostatic stress” reflects the stress states of the TiC particles. Irregular fluctuations are observed in Al6061+5%TiC under stress values of −75 to −120 MPa (Figure 6i), showing a complex hydrostatic stress change. This agrees with the possible interphase stress changes due to the initial plastic deformation in Al near the phase boundary, aligning with the second stage of deviatoric stress evolution.

3.4. Effects of TiC Particles on Dislocation Activity at 250 °C

Adding TiC not only affects load sharing but also alters the deformation behavior of the Al matrix in terms of dislocation activity and texture development. To assess texture evolution, changes in the peak intensities of different Al lattice planes in Al6061 and Al6061+TiC samples during loading at room temperature and 250 °C were quantified and compared (Figure 7).

At room temperature, the peak intensity of Al {220} increases, while those of Al {111} and Al {200} decrease (Figure 7a–c). This pattern is typical for aluminum alloys loaded at room temperature, in which only the {111} slip plane is activated. At elevated temperatures, the slip on the {101} plane is activated as well, in addition to that on the {111} plane [33]. The new slip system is thought to result in an increase in {200} peak intensity along the compression direction, which is the opposite of the result of the slip system at room temperature. The competition of these two systems mitigates the changes in Al {200} intensity during the deformation of as-printed Al6061 at 250 °C. Specifically, at a strain of 5%, the peak intensity of Al {200} decreases by 1.6% at 250 °C, compared with a 13.6% decrease at room temperature in Al6061 (Figure 7d). Interestingly, this competition can be altered by adding TiC particles. Figure 7e,f show a further increase in the peak intensity of Al {200} in the composites, with a 1.5% increase at a strain of 5% for Al6061+5%TiC. This suggests that TiC may inhibit the dislocation slip on the {111} plane more than that on the {101} plane at 250 °C.

Adding TiC particles also affects the amplitude of the peak intensities. At both room temperature and 250 °C, the amplitude of the peak intensity change decreases as the volume fraction of TiC increases from 2% to 5%. This effect likely occurs because more TiC particles hinder dislocations around them to strengthen the matrix instead of allowing dislocations to slip through the whole grain and develop deformation textures.

While they hinder the dislocation movement, TiC particles contribute to the accumulated dislocations in the Al matrix during the deformation, overcoming the annihilation of dislocations at 250 °C. As the higher dislocation density broadens the diffraction peaks, the peak width of different Al lattice planes during loading at 250 °C is quantified and plotted in Figure 8. In Al6061, the peak width shows an overall decreasing trend after a slight initial increase (Figure 8a), indicating that dislocation generally exhibits annihilation [58]. By contrast, the peak width in the Al6061+TiC samples shows a continuous increase after an initial rapid rise (Figure 8b,c), suggesting overall dislocation multiplication with the presence of TiC. This difference arises because TiC particles can pin dislocations in the Al matrix, promoting dislocation accumulation around the TiC particles. This accumulation counteracts dislocation annihilation, resulting in an overall increase in dislocation density in Al6061+TiC. This also explains the continuous hardening observed in the Al6061+TiC samples, in contrast to the softening observed in Al6061 (Figure 2).

4. Conclusions

In this work, Al6061+TiC composites were additively fabricated by laser powder bed fusion (LPBF) and used as a model material for in situ neutron diffraction studies to elucidate phase-specific deformation behaviors at an elevated temperature. In situ neutron diffraction was demonstrated to capture the structural evolution of the Al matrix, dispersive TiC phase, and minor precipitate from the alloy during the compression test at 250 °C. It was found that the addition of a small amount of nano-size TiC significantly increases the strength and alters the deformation behavior at 250 °C. The deviatoric stresses calculation excluded the influence of chemistry changes on the stress analysis, and it revealed the three load-sharing stages in the phases. Further analysis of Bragg peak intensity and broadening reveals that the presence of TiC altered the dislocation activity during deformation at 250 °C by influencing dislocation slip planes and promoting dislocation accumulation. The key findings are summarized below:

(1)

Adding TiC significantly enhances yield strength, which increases from 66 MPa for Al6061 to 96 MPa for Al6061+2%TiC and further to 104 MPa for Al6061+5%TiC at 250 °C.

(2)

Unlike the two-stage elastic–plastic behavior observed in Al6061, Al6061+TiC composites exhibit three stages during compression: (1) the elastic deformation of both the Al and TiC phases; (2) the early yielding of the Al matrix triggered by local plastic deformation near TiC particles, and (3) the global plastic deformation of the Al matrix.

(3)

Minor chemical changes due to Mg₂Si precipitation and the dissolution of Al into TiC were identified. These changes lead to fluctuations in the measured “phase stresses” during loading at 250 °C. The addition of TiC reduces the deviatoric stress carried by the Al matrix, delaying the yielding of the metal matrix and thereby enhancing the composite’s strength.

(4)

Peak intensity quantification revealed that in addition to the {111} slip plane typically activated at room temperature, the {101} slip system is also activated at 250 °C. The addition of TiC inhibits dislocation slip on the {111} plane more than on the {101} plane at 250 °C.

(5)

The peak width of different Al lattice planes shows an overall increasing trend in Al6061+TiC, contrasting with the overall decreasing trend in pure Al6061 at 250 °C. This difference arises because TiC particles pin dislocations in the Al matrix, promoting dislocation accumulation around TiC particles and thereby increasing dislocation density in Al6061+TiC.

These findings provide direct experimental observations of the phase-specific dynamic process in AMCs under deformation at an elevated temperature. The revealed mechanisms provide insights for the future design and optimization of high-performance AMCs.