In cold rolling of a thin strip, the contact angle *α* is very small; thus, cos *α* ≈ 1. The mass flow relationship in the deformation zone can be expressed as *hxvx* = *h*1*v*1 = *h*2*v*2. Substituting Equations (8)–(10) into Equation (26), we obtain

$$\begin{aligned} A\_{l1} &= \frac{Klf\_{1}}{\delta\_{1}\Delta h} \left[ \frac{h\_{2}v\_{2}(\delta\_{1}\varepsilon\_{2}+1)\left(\mu\_{1}^{\delta\_{1}-1}-1\right)}{\delta\_{1}} - h\_{2}v\_{2}\ln\mu\_{l} - \frac{hv\_{1}(\delta\_{1}\varepsilon\_{2}+1)\left(\mu\_{1}^{\delta\_{1}+1}-1\right)}{\delta\_{1}+1} + v\_{1}(h\_{1}-h) \right] \\ &+ \frac{Klf\_{1}}{\delta\_{2}\Delta h} \left\{ \frac{h\_{2}v\_{2}\left(\mu\_{a}^{\delta\_{2}-1}-1\right)}{\delta\_{2}-1} \left[\frac{\delta\_{2}}{\delta\_{1}}(\delta\_{1}\varepsilon\_{1}-1)\mu\_{b}^{\delta\_{1}}+\frac{\delta\_{2}}{\delta\_{1}}-1\right] + v\_{1}(h\_{2}-h\_{1}) \right. \\\\ &\left. - \frac{h\_{2}v\_{2}\left(\mu\_{a}^{\delta\_{2}-1}-1\right)}{\delta\_{2}} \left[\frac{\delta\_{2}}{\delta\_{1}}(\delta\_{1}\varepsilon\_{1}-1)\mu\_{b}^{\delta\_{1}}+\frac{\delta\_{2}}{\delta\_{1}}-1\right] - h\_{2}v\_{2}\ln\mu\_{c} \right. \\\\ &\left. + \frac{Kf\_{1}}{\delta\_{1}\Delta h} \left[\frac{h\upsilon\_{1}(\delta\_{1}\varepsilon\_{1}-1)\left(\mu\_{b}^{\delta\_{1}-1}-1\right)}{\delta\_{1}-1} + v\_{1}(H-h\_{2}) - \frac{h\upsilon\_{2}v\_{2}(\delta\_{1}\varepsilon\_{1}-1)\left(\mu\_{b}^{\delta\_{1}}-1\right)}{\delta\_{1}} \right. \\\\ &\left. - h\upsilon\_{2}\upsilon\_{2}\ln\mu\_{$$

Friction power on the lower surface of the strip in the deformation zone can be obtained using the same method.

$$\begin{aligned} A\_{l2} &= \frac{Klf\_{2}}{\delta\_{1}\Delta h} \left[ \frac{h\_{2}v\_{2}(\delta\_{l}\varepsilon\_{2}+1)\left(\mu\_{l}^{\delta\_{1}-1}-1\right)}{\delta\_{1}} - h\_{2}v\_{2}\ln\mu\_{l} - \frac{hv\_{2}(\delta\_{l}\varepsilon\_{2}+1)\left(\mu\_{l}^{\delta\_{1}+1}-1\right)}{\delta\_{1}+1} + v\_{2}(h\_{1}-h) \right] \\ &+ \frac{Klf\_{2}}{\delta\_{2}\Delta h} \left\{ -\frac{hv\_{2}v\_{1}\left(\mu\_{l}^{\delta\_{2}-1}-1\right)}{\delta\_{2}-1} \left[\frac{\delta\_{2}}{\delta\_{1}}\left(\delta\_{1}\varepsilon\_{1}-1\right)\mu\_{b}^{\delta\_{1}} + \frac{\delta\_{2}}{\delta\_{1}} - 1\right] + v\_{2}(h\_{2}-h\_{1}) \right. \\\\ &+ \frac{h\_{2}v\_{2}\left(\mu\_{l}^{\delta\_{2}-1}-1\right)}{\delta\_{2}} \left[\frac{\delta\_{2}}{\delta\_{1}}\left(\delta\_{1}\varepsilon\_{1}-1\right)\mu\_{b}^{\delta\_{1}} + \frac{\delta\_{2}}{\delta\_{1}} - 1\right] - h\_{2}v\_{2}\ln\mu\_{c} \right. \\\\ &+ \frac{Klf\_{2}}{\delta\_{1}\Delta h} \left[\frac{hv\_{2}\left(\delta\_{1}\varepsilon\_{1}-1\right)\left(\mu\_{b}^{\delta\_{1}-1}-1\right)}{\delta\_{1}-1} + v\_{2}(H-h\_{2}) - \frac{h\_{2}v\_{2}\left(\delta\_{1}\varepsilon\_{1}-1\right)\left(\mu\_{b}^{\delta\_{1}}-1\right)}{\delta\_{1}} \right. \\\\ &- h\_{2}v\_$$

The friction power per unit width can be expressed as *A*f = *A*f1 + *A*f2.

#### *2.8. Energy Consumption in Asymmetrical Rolling*

The energy consumption in asymmetrical strip rolling refers to the work consumed during the process of rolling a strip of a certain size to target thickness. In cold strip manufacturing, the main energy-consuming process is rolling. Therefore, only energy consumption in the rolling process of asymmetrical strip rolling is investigated in this work. The energy consumption in the rolling process mainly consists of energy consumption in rolling, energy consumption of auxiliary equipment and energy consumption due to friction in the drive system. To simplify the analysis, only the energy consumption in rolling is taken into account in this work.

The energy consumption in rolling is equivalent to the work completed by the motors of the rolling mill and the motors of the coiler and uncoiler. The energy consumption in one pass can be expressed as

$$\mathcal{W}\_{\rm T} = \mathcal{W}\_1 + \mathcal{W}\_2 + \mathcal{W}\_{\rm rf} + \mathcal{W}\_{\rm rb} \tag{29}$$

where *W*1 = *A*1*B*0*t* 60000 , *W*2 = *A*2*B*0*t* 60000 , *W*rf = *A*rf*B*0*t* 60000 , *W*rb = *A*rb*B*0*t* 60000 .

The work performed in the deformation during plastic working is equal to the work of external forces. In a rolling pass, rolling work (*W*T) is equivalent to the sum of the plastic deformation work of the strip (*W*d) and the friction work on the contact surface between the strip and rolls (*W*f).

$$\mathcal{W}\_{\mathbb{T}} = \mathcal{W}\_{\mathbb{d}} + \mathcal{W}\_{\mathbb{f}} \tag{50}$$

Friction work *W*f consists of *W*f1 (on the upper surface) and *W*f2 (on the lower surface)

$$\mathcal{W}\_{\rm f} = \mathcal{W}\_{\rm f1} + \mathcal{W}\_{\rm f2} \tag{31}$$

where *W*f1 = *A*f1*B*0*t* 60000 , *W*f2 = *A*f2*B*0*t* 60000 .

The proportion of friction work (*Q*f) and proportion of plastic deformation work (*Q*d) can be expressed as *Q*f = *Q*f *W*T × 100% and *Q*d = *Q*d *W*T × 100%, respectively.

#### **3. Results and Discussion**

Several asymmetrical strip rolling experiments have been conducted on the asymmetrical mill, with a work roll diameter of 90 mm, to verify the validity of the proposed models. The lower roll served as the slow roll and was set to a peripheral velocity of 4.0 m/min, while the peripheral velocity of the upper roll varied with the speed ratio. Emulsion was used as a lubricant. The 430 stainless steel strips, with 80.0 mm in width, were rolled from 0.150 mm to 0.105 mm under different speed ratios. By fitting the tensile test data of 430 stainless steel strips with different reduction ratios, the plane deformation resistance can be expressed as

$$K = 407.9 + 588.6\overline{\varepsilon}^{0.718} \tag{32}$$

where *ε* = 0.4 1 − *H H*0 + 0.6 1 − *h H*0 .

The model-calculated roll force is compared with the experiment-measured roll force. A comparison of roll force between experimental measurement values and those calculated by the models is illustrated in Figure 3a. It was found that both the experimental and the calculated roll force decreased with an increasing speed ratio. The maximum error and RMS error between the experimentally measured roll forces and calculated ones are 9.7% and 5.9%, respectively. The model-calculated forces are also compared with the results of models from the literature in Figure 3b. Models from the literature cannot finish calculations in the third point. All model-calculated roll forces are less than experimental values, but the present model has better accuracy than models from the literature. Data used in the calculation are listed in Table 1.

The proposed models are used to calculate rolling pressure, roll force, roll torque and roll power, as well as friction work and plastic deformation work in asymmetrical strip rolling. A work roll diameter of 88 mm, and 430 stainless steel strips of initial length *L*0 = 1000 m and initial width *B*0 = 100 mm, are used in the analysis. From the calculation results, the effects of rolling parameters on friction work and plastic deformation work are analyzed.

**Figure 3.** Comparisons of model-calculated roll force with: (**a**) experimental results; (**b**) experimental results and results of literature models.

**Table 1.** Data for roll force calculation.

Figure 4 illustrates the effects of the speed ratio on friction work, plastic deformation work and rolling work in a single rolling pass with the same reduction ratio. It can be seen that the upper friction work increases significantly with an increasing speed ratio, but the lower friction work only increases slightly. This is because the percentage of cross-shear zone increases with an increasing speed ratio. As a result, the upper and lower neutral points move toward the exit and the entrance of the deformation zone, respectively, thus increasing the slipping distance between the upper roll and the strip. Therefore, the friction work increases. The plastic deformation work remains unchanged under a different speed ratio because the deformation degree is the same.

**Figure 4.** Effects of speed ratio on: (**a**) friction work, (**b**) plastic deformation work and rolling work, in a single pass.

Figure 5 illustrates the effects of the speed ratio on total friction work, total plastic deformation work, total rolling work, the proportion of total friction work and the proportion of total plastic deformation work in multi-pass rolling. The curves in Figure 5a,b were obtained from five-pass rolling with the same reduction ratio, while the curves in Figure 5c,d were obtained from multi-pass rolling with the same roll force. It is seen that both the total friction work and the total rolling work increase as the speed ratio increases, but the total plastic deformation work remains constant, as in Figure 5a,c. Owing to the same level of accumulated reduction, the total plastic deformation work is not affected by the speed ratio. The sum of accumulated relative velocity between rolls and strips and the sum of the slipping distance increase with an increasing speed ratio. Consequently, the total friction work increases, leading to an increase in the total rolling work. Correspondingly, the proportion of total friction work gradually increases, and the proportion of total plastic deformation work gradually decreases with an increasing speed ratio, as shown in Figure 5b,d. The proportion of total friction work exceeds the proportion of total plastic deformation work at a certain speed ratio.

**Figure 5.** Sum of works and proportions of friction work and plastic deformation work in multi-pass rolling: Work (**a**), proportion of friction work and plastic deformation work (**b**) in a five-pass rolling with same reduction ratio; work (**c**), proportion of friction work and plastic deformation work (**d**) in multi-pass rolling with same roll force.

Figure 6 illustrates the effects of the speed ratio and entry thickness on friction work, plastic deformation work and its proportions. According to Figure 6, we know that the plastic deformation work increases with both an increasing speed ratio and an increasing entry thickness. This is because both an increase in entry thickness and an increase in speed ratio increase the magnitude of deformation, thus resulting in an increase in plastic deformation work. Friction work increases with an increasing speed ratio for the same entry thickness, and also increases with increasing thickness for the same speed ratio. As the speed ratio increases, the percentage of cross-shear zone increases, and accumulated

relative velocity between rolls and strips consequently increases. As a result, friction work increases.

**Figure 6.** Variation of (**a**) plastic deformation work, (**b**) friction work, and (**c**) proportions of plastic deformation work and friction work with entry thickness.

Figure 7 illustrates the effects of the friction coefficient and speed ratio on total friction work and its proportions in total rolling work. As the speed ratio increases, both the total friction work and the proportion of total friction work increase. The total friction work and the proportion of total friction work also increase with an increased friction coefficient.

**Figure 7.** Variation of (**a**) total friction work and (**b**) proportion of total friction work with friction coefficient.

The curves in Figure 8 show the effects of plane deformation resistance and the speed ratio on the total plastic deformation work, total friction work and proportion of total friction work in multi-pass rolling. Given that the deformation force increases with an increase in plane deformation resistance, and the levels of accumulated reduction are the same, the total plastic deformation work increases with an increase in plane deformation resistance. The increase in plane deformation resistance causes the percentage of the cross-shear zone to increase, thus increasing the total friction work. The proportion of total friction work decreases with an increase in plane deformation resistance because the increment rate in the total plastic deformation work is greater than in the total friction work.

**Figure 8.** Variation of (**a**) total plastic deformation work, (**b**) total friction work, and (**c**) proportion of total friction work with plane deformation resistance.

Figure 9 shows the effects of the speed ratio on the total friction work and the proportions of total friction work under various back tensions. The percentage of the cross-shear zone decreases with the increase in back tension, resulting in a decrease in the relative slip between the strip and roll. Therefore, the total friction work decreases slightly with an increasing back tension.

The effects of front tension are similar to that of back tension.

The effects of the entry thickness and the friction coefficient on the proportions of friction work are shown in Figure 10. There is a special line on which the proportion of friction work is equal to 50%; in other words, the proportion of friction work is the same as the proportion of plastic deformation work. At points T1, T2, and T3 in Figure 10a, which are called the threshold points, the curve is divided into two parts, and the friction work is greater than the plastic deformation work on the left part. On the contrary, at points S1, S2, and S3 in Figure 10b, the friction work is greater than the plastic deformation work on the right part.

**Figure 9.** Variation of (**a**) total friction work and (**b**) proportion of total friction work with back tension.

**Figure 10.** Variation of threshold point and proportions of friction work with: (**a**) entry thickness; (**b**) friction coefficient.

The threshold point is related to friction conditions in the deformation zone and the deformation degree; it is affected by parameters such as thickness, the friction coefficient, the speed ratio, front and back tension, deformation resistance and the reduction ratio.

#### **4. Conclusions**

Mathematic models based on the slab method are built to calculate the rolling pressure, roll force, roll torque, friction power, rolling power and rolling work for the three typical region deformation zone configurations in the asymmetrical rolling of a thin strip. The effects of rolling parameters on energy consumption (friction work and plastic deformation work) are investigated by analyzing computed results. The following conclusions are obtained.

**Author Contributions:** Conceptualization, X.S.; methodology, X.S.; software, Q.Z. and X.S.; validation, Q.Z., X.S. and X.L.; formal analysis, Q.Z. and X.S.; investigation, Q.Z.; resources, X.S.; data curation, Q.Z.; writing—original draft preparation, Q.Z.; writing—review and editing, X.L.; visualization, Q.Z.; supervision, X.L.; project administration, X.L. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was financially supported by the program for the National Natural Science Foundation of China No. 51374069.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

## *Article* **Effect of CeO2 Size on Microstructure, Synthesis Mechanism and Refining Performance of Al-Ti-C Alloy**

**Yanli Ma 1, Taili Chen 2, Lumin Gou 2 and Wanwu Ding 1,2,\***

1 Technology Center, Jiuquan Iron and Steel (Group) Co., Ltd., Jiayuguan 735100, China<EMAIL_ADDRESS>2 School of Materials Science and Engineering, Lanzhou University of Technology, Lanzhou 730050, China;

<EMAIL_ADDRESS>(T.C<EMAIL_ADDRESS>(L.G.)

**\*** Correspondence<EMAIL_ADDRESS>Tel.: +86-1511-7183-365

**Abstract:** The effects of CeO2 size on the microstructure and synthesis mechanism of Al-Ti-C alloy were investigated using a quenching experiment method. A scanning calorimetry experiment was used to investigate the synthesis mechanism of TiC, the aluminum melt in situ reaction was carried out to synthesize master alloys and its refining performance was estimated. The results show that the Al-Ti-C-Ce system is mainly composed of α-Al, Al3Ti, TiC and Ti2Al20Ce. The addition of CeO2 obviously speeds up the progress of the reaction, reduces the size of Al3Ti and TiC and lowers the formation temperature of second-phase particles. When the size of CeO2 is 2–4 μm, the promotion effect on the system is most obvious. The smaller the size of CeO2, the smaller the size of Al3Ti and TiC and the lower the formation temperature. Al-Ti-C-Ce master alloy has a significant refinement effect on commercial pure aluminum. When the CeO2 size is 2–4 μm, the grain size of commercial pure aluminum is refined to 227 μm by Al-Ti-C-Ce master alloy.

**Keywords:** CeO2 size; quenching experiment method; DSC; refining performance

**Citation:** Ma, Y.; Chen, T.; Gou, L.; Ding, W. Effect of CeO2 Size on Microstructure, Synthesis Mechanism and Refining Performance of Al-Ti-C Alloy. *Materials* **2021**, *14*, 6739. https://doi.org/10.3390/ma14226739

Academic Editor: Daniela Kovacheva

Received: 6 September 2021 Accepted: 2 November 2021 Published: 9 November 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

#### **1. Introduction**

A number of studies show that grain refiner for aluminum and its alloys can significantly improve the mechanical properties, casting properties, deformation processing properties and surface quality of materials [1,2].

Al-Ti-B master alloy has been the most widely used grain refiner and has been applied commercially [3]. However, many defects have been found in the application of Al-Ti-B master alloy thus limiting its development, i.e., a larger aggregation tendency of TiB2 in the melt, large size, and the possibility of Zr and Si poisoning [4,5]. In order to obtain a good refining efficiency, a series of novel master alloys have emerged, such as Al-Ti-C, Al-Ti-B-C and Al-Ti-C/B-RE master alloy [6–9]. Al-Ti-C master alloy especially has been extensively studied by researchers, and it has been found that it can be remedy certain defects of Al-Ti-B master alloy [10,11]. Unfortunately, the refining efficiency of Al-Ti-C is unstable, and the possible reasons for this are that the active of TiC is insufficient, TiC in the melt aggregates easily and the morphology of Al3Ti and TiC is easily affected by preparation conditions. For the most part, the wetting between Al and C is poor, and the formation of TiC is therefore difficult [12].

Ao et al. [13] found that La2O3 added to the Al-Ti-C-CuO system could promote the wettability of C and Al melt. The results of thermodynamic analysis from differential scanning calorimetry (DSC) experiments showed that La2O3 could promote metastable phase transfer into stable phases. Furthermore, the even distribution of TiC can directly influence the refining efficiency of Al-Ti-C. Wang L. D. et al. [14] analyzed the thermodynamics of the Al-Ti-C-Ce system using DSC experiments, and the result showed that rare earth oxide promoted the reaction and made the distribution of TiC even, forming a new Ti2Al20Ce phase. Moreover, CeO2 has a catalytic effect and can promote the reaction. Korotcenkov et al. [15] considered that the catalytic activity of CeO2 is closely related to its structure, morphology

and size. The research of Auffan et al. [16] and Liu et al. [17] shows that CeO2 nanoparticles have a high specific surface area, and the smaller the size of CeO2 nanoparticles, the higher the specific surface area and the greater the activity. Preliminary studies have shown that when the content of CeO2 was 4 wt.%, it had a significant promoting effect on the Al-Ti-C system [18]. On this basis, it is necessary to further study the effect of CeO2 particle size on phase transition and microstructure transformation of Al-Ti-C system.

In this study, the aluminum melt in situ reaction was carried out to prepare Al-Ti-C-Ce master alloys. The influence of the size of CeO2 on the thermodynamic process, reaction products and microstructures of the prepared Al-Ti-C-Ce system were studied in detail through quenching experiments, and the reaction mechanism of the Al-Ti-C-Ce system is summarized in this paper. The goal of the abovementioned studies was to determine the influence mechanism on the microstructure and phase transformation with different sizes of CeO2 in the Al-Ti-C-Ce system based on thermodynamic analysis and DSC experiments.

#### **2. Experiment**

#### *2.1. Quenching Experiments*

Al powders (99.0 wt.%, 80–100 μm), Ti powders (99.0 wt.%, 45–65 μm) and C powders (99.0 wt.%, 10–20 μm) were prepared with Al/Ti/C = 5:2:1 (molar ratio) and 4 wt.% CeO2 powder (99.0 wt.%, with different average sizes, i.e., 30 nm, 50 nm, 1 μm, 2–4 μm). The mixture was evenly mixed in a PULVERISETTE-5 high-speed planetary ball mill with a ballto-material ratio = 3:1. The rotation speed and the total milling time were 350 r/min and 3 h, respectively. Then, 50 g of the mixture was cold-pressed into a cylindrical prefabricated block of Φ 25 mm × 50 mm on an AG-10TA universal test stretching machine (Shimadzu Corporation, Kyoto, Japan) and dried in a drying oven at 200 ◦C for 2 h.

Simultaneously, commercial pure aluminum (99.7 wt.%, A 99.7) was melted in alumina crucible by using a SG-7.5–10 type crucible furnace (Zhonghuan Experimental Furnace Corporation, Tianjin, China). After the temperature was raised to 800 ◦C, the prefabricated block was added to the melt through the graphite bell jar. According to the previously determined typical sampling time (8, 50, 60, 80 and 90 s), the reaction block was quickly removed at each time point and subjected to high pressure ice brine flow quenching to obtain quenched samples.

The specimens with a size of 10 mm3 taken from the quenched samples were electrolytically polished by a reagent (10 vol.% HClO4 + 90 vol.% absolute ethanol, voltage is 20 V and the time is about 10~30 s) after mechanical grinding and polishing. The phase constituents of specimens were analyzed using a D8 advance X-ray diffractometer (XRD, Shimadzu Corporation, Kyoto, Japan). The microstructure was examined using a JSM-6700F scanning electron microscope (SEM) equipped with an energy dispersive spectrometer (EDS, Shimadzu Corporation, Kyoto, Japan). In addition, DSC analysis of the Al-Ti-C-Ce system was conducted with a NETZSCH STA 449F3 instrument (NETZSCH, Hanau, Germany) with a heating rate of 20 K/min and samples weighing less than 10 mg.

#### *2.2. Grain Refinement Experiments*

To evaluate the grain refining effect of Al-Ti-C master alloy after adding CeO2, the different master alloys were prepared using the aluminum melt in situ reaction method, and then master alloys were used to refine A 99.7. The specific experimental process is as follows:

Firstly, A 99.7 ingots were melted and heated up to 800 ◦C in an alumina crucible by using a SG-7.5–10 type crucible furnace (Zhonghuan Experimental Furnace Corporation, Tianjin, China). A proper quality of prefabricated block was then added to the melt after preheating at 200 ◦C for 2 h. After being held at 800 ◦C for 15 min, the melt was finally poured into a steel die with an inner diameter of 45 mm, a diameter of 70 mm and a height of 70 mm. By adjusting the size and addition of CeO2, Al-Ti-C and Al-Ti-C-Ce master alloys were obtained. The specific experimental parameters and numbers of master alloys are shown in Table 1.

**Table 1.** The experimental parameters and numbers of master alloys.

Thereafter, an appropriate amount of A 99.7 was melted in a crucible electrical resistance furnace at about 730 ◦C. The 0.3 wt.% master alloys were sectioned at the position of 10 mm from the bottom on #1~#5 alloys and then added to the melt. After that, the melt was held for 10 min and finally poured into the steel mold. By solidification, the refined samples were obtained. For comparison, standard A 99.7 was treated in the same process. A specimen with a thickness of 10 mm was cut along the bottom of the refined sample. A cross profile of the specimen was etched using Keller's reagent (25 vol.% H2O + 15 vol.% HNO3 + 15 vol.% HF + 45 vol.% HCl). The grain structures of specimens were taken by a camera, and the average grain sizes were calculated using the linear intercept method.

#### **3. Results and Discussion**

#### *3.1. The Raw Materials of CeO2 and the Mixed Particles*

The raw materials of different CeO2 size are presented in Figure 1. From Figure 1a,b, it can be seen that the morphology of CeO2 with a size of 30 and 50 nm is spherical nanoparticles. In contrast, the 1 and 2–4 μm sized CeO2 gathered together in a short rod shape and small block shape, as shown in Figure 1c,d.

**Figure 1.** Different sizes of CeO2: (**a**) 30nm; (**b**) 50nm; (**c**) 1 μm; (**d**) 2–4 μm.

Figures 2 and 3 show the microstructure and EDS results of the mixed powders (Al, Ti, C and CeO2) with different CeO2 size after ball milling. As shown Figure 2, Al, Ti, C and CeO2 powders were evenly mixed after ball milling. Combining the EDS mapping analysis of Al, Ti, C, Ce and O elements as shown in Figure 3, it can be seen that the bright white particles with larger size are Ti, gray particles are Al and black particles are C. The bright white particles with small size are mainly enriched with Ce and O elements, which can be judged as CeO2 particles. As can be observed in Figure 2a,b, for nano-CeO2 particles, ball milling mainly leads to uniform dispersion. However, for the mixed powder of CeO2 with a larger size, ball milling not only contributes to dispersion but also destroys the aggregation between CeO2, breaking the CeO2 into smaller particles. A statistical analysis of the particle size of CeO2 found that after ball milling, the size of 1 μm CeO2 particles was reduced to about 400 nm (see Figure 2c), and the size of 2–4 μm CeO2 particles was reduced to about 200 nm (see Figure 2d).

**Figure 2.** The microstructure of the mixed powders (Al, Ti, C and CeO2) with different CeO2 size after ball milling: (**a**) 30 nm; (**b**) 50 nm; (**c**) 1 μm; (**d**) 2–4 μm.

#### *3.2. Effect of CeO2 Particle Size on Phase Transformation and Microstructure Transformation of Al-Ti-C Master Alloy during Preparation*

In order to study the influence of CeO2 size on the microstructure transformation process of Al-Ti-C system, the microstructure of quenched samples with different CeO2 sizes at the same reaction time (60 s) was analyzed. SEM images and EDS results are shown in Figure 4. Comparing Figure 4a,b, it is found that when the CeO2 size is 30 and 50 nm, it has little effect on the reaction process of the system. As can be seen, the molten Al closely connects Ti particles, and incomplete Ti particles and incomplete melting Al particles were reserved in system after solidification. As can be seen from Figure 4c, a variety of compounds with different sizes and morphologies are formed around the Ti particles, which are Al3Ti, TiC, Al-Ti compounds and CeC2 particles. It can be seen from Figure 4d that when the CeO2 size is 1 μm, there is no close connection between the Ti particles and the aluminum matrix. Otherwise, compared to Figure 4a,b the degree of reaction is low, and there are obvious incompletely reacted Al, Ti and C particles in the system. When CeO2 with a size of 2–4 μm is added to the system, as shown in Figure 4e,f, it can be seen that the size of the Ti particles is reduced, the reaction intensity is enhanced, the density of the system is increased and the particles are closely connected. Especially, A large number of TiC, Al3Ti and Ti2Al20Ce phases are formed in the Al matrix.

**Figure 3.** SEM image and map scanning patterns of Area 1 in Figure 2d: (**a**) SEM image; (**b**–**f**) map scanning patterns of Al, Ti, C, Ce and O elements, respectively.

**Figure 4.** The microstructure of quenching samples with different CeO2 size at 60 s: (**a**) 30 nm; (**b**) 50 nm; (**c**) SEM image of Area 1; (**d**) 1 μm; (**e**) low magnification image and (**f**) high magnification image of samples with CeO2 of the size of 2–4 μm.

This difference in response is because the larger the CeO2 particles, the larger the specific surface area and the stronger the catalytic effect, which obviously accelerates the progress of the reaction. For micron-sized CeO2 particles, due to the crushing of CeO2 particles during ball milling, more active sites are provided, and the larger the particle size, the more severe the degree of crushing, the more active sites and the greater the promotion of the reaction.

The microstructure of the quenched samples with different CeO2 sizes added under complete reaction was observed. The phase, microstructure characteristics and the average grain size of the second phase particles of the quenched samples were analyzed, as shown in Figures 5–7.

**Figure 5.** XRD patterns of prefabricated blocks with different CeO2 size under complete reaction.

Figure 5 shows XRD patterns of prefabricated blocks with different CeO2 size under complete reaction. As can be seen from Figure 5, under complete reaction, the main phases in the sample are α-Al, Al3Ti, TiC and Ti2Al20Ce. When the CeO2 particle size is 30 nm and 2–4 μm, the diffraction peak intensity of TiC (2θ = 41.7◦) is higher than 50 nm and 1 μm, and the diffraction peak intensity of Al3Ti, TiC and Ti2Al20Ce showed a weakening trend with the increase in CeO2 size.

Figure 6 shows SEM images of the quenching samples of prefabricated block with different CeO2 size under complete reaction. Comparing Figure 6a–d, it is found that the size of CeO2 particles has a great influence on the quantity, morphology and distribution of Al3Ti and TiC. When the size of CeO2 is 30 nm and 2–4 μm, the amount of TiC and Al3Ti is large and the size is relatively uniform, the distribution is dispersed in the matrix. When the size of CeO2 is 50 nm, TiC and Al3Ti adhere to each other in an irregular shape, and the size is not uniform. When the size is 1 μm, the number of TiC particles is small, and Al3Ti is evenly distributed in the matrix.

Figure 7 shows the statistics chart of the second phase particles in prefabricated block with different CeO2 size under complete reaction. It can be seen from Figure 7 that the particle size of CeO2 mainly affects the size of TiC particles, and has little effect on the size of Al3Ti. When the size of CeO2 is 30 nm, the average size of Al3Ti and TiC is the smallest at 4.0 and 1.1 μm, respectively. With the increase in CeO2 particle size, the average size of TiC gradually increases, the average size of Al3Ti increases first and then tends to be flat and CeO2 particle size is 2–4 μm. The average size of Al3Ti in the quenched samples is 6 μm, and the average size of TiC is about 2.2 μm.

**Figure 6.** SEM images of the quenching samples of prefabricated block with different CeO2 size under complete reaction: (**a**) 30 nm; (**b**) 50 nm; (**c**) 1 μm and (**d**) 2–4 μm.

**Figure 7.** The statistics chart of the second phase particles in prefabricated block with different CeO2 size under complete reaction.

From the above analysis, it can be seen that the content and size of CeO2 have a significant effect on the formation, distribution, size, number and morphology of phases in the Al-Ti-C system. In order to further study the formation and transformation process of various phases, the pre-fabricated blocks with CeO2 size of 2–4 μm and 4 wt.% were taken as the research objects to study the reaction products and microstructures of the Al-Ti-C-Ce system at different stages. The XRD pattern and microstructure are shown in Figures 8–10.

**Figure 8.** XRD spectra of prefabricated blocks with CeO2 particle size of 2–4 μm at different reaction times: (**a**) 8, (**b**) 50, (**c**) 60, (**d**) 80, (**e**) 90 s.

**Figure 9.** SEM images of #5 quenching samples prepared at different times: (**a**) 8; (**b**) 50; (**c**) 60; (**d**) 80 and (**e**) 90 s.

**Figure 10.** The map scan patterns of Figure 9d: (**a**) backscattered electron image of Figure 9d; (**b**–**e**) map scanning patterns of Al, Ti, C and Ce elements, respectively.

As can be seen from Figure 8a, in the initial stage of the reaction (8 s), the main detection phases in the system are Ti, C, CeO2 and a small amount of Al1+xTi1-x, indicating that there is a weak solid–solid diffusion process between solid Al and solid Ti at this stage. When the reaction time reaches 50 s, TiC, Al3Ti and trace Ti2Al20Ce are formed, as shown in Figure 8b. As the reaction progresses, the intensity of the diffraction peaks of Al3Ti and TiC decreases first and then increases. The intensity of the diffraction peaks of Ti2Al20Ce continues to increase. When the reaction reaches 90 s, a large amount of TiC, Al3Ti and Ti2Al20Ce are synthesized in the system; only weak incomplete reaction Ti peaks were detected in Figure 8c–e.

As can be seen from Figure 9a, the system basically maintains the original particle state. Figure 9b shows that when the reaction reaches 50 s, the molten Al in the system spreads on the surface of the Ti particles, and a small amount of Al1+xTi1-x is formed under the solid–liquid diffusion reaction [19,20] and then transformed into Ti2Al20Ce under the action. In addition, around 3 μm Al3Ti and 1 μm TiC were formed around the Ti particles, and unreacted Ti particles and CeO2 also existed in the system. Figure 9c shows that there is a size of about 5 μm Al3Ti, and bulk Ti2Al20Ce are speculated to be caused by the enrichment of the Al element and the Ce element around the Ti element. In addition, solid C reacts with Ti to form a large number of TiC particles, and the whole system reacts to form a wrapped structure of Ti/Al-Ti/Al3Ti/Ti2Al20Ce [21,22]. The results of Figure 9d and its surface scanning energy spectrum (see Figure 10) show that rare earth elements are mainly enriched with the surface of C element and around the Ti element, and some rare earth elements react with dissolved Ti atoms and liquid Al to form Al3Ti and Ti2Al20Ce. The other part reacts with solid C to form CeC2 and then reacts with dissolved Ti atoms to generate TiC. At the same time, because the affinity between Ti atoms and C atoms is greater than that of Al and C, dissolved Ti atoms directly generate TiC particles with dissolved C atoms. As time goes on, TiC particles grow up, and the encapsulated structure gradually disappears. Al3Ti particles are free from the encapsulated structure and freely distributed in the matrix, and the rare earth Ce element continues to enrich around Al3Ti, which hinders the growth of Al3Ti particles. The aspect reacts with it to generate Ti2Al20Ce and grows up by combining [22,23], as shown in Figure 9e. In Figure 9e, TiC particles with an average size of about 2 μm, small bulk Al3Ti with an average size of about 6 μm and

irregular bulk Ti2Al20Ce with an average size of about 35 μm were formed in the late stage of the reaction.

In order to further explore the specific reaction process of the Al-Ti-C-Ce system, the influence of CeO2 on the TiC synthesis method of Al-Ti-C system was studied. Combined with the DSC curve of the Al-Ti-C system in the previous research report [24], differential scanning calorimetry analysis was performed on the preforms with CeO2 size of 30 nm and 2–4 μm. The DSC curve is shown in Figure 11. The corresponding reaction heat changes are shown in Table 2. From the curve, an obvious endothermic peak and two exothermic peaks can be seen. The first endothermic peaks at temperatures of 670.5 and 671.2 ◦C represented the melting of aluminum, and the corresponding reaction endotherms were 173.5 and 195.5 J/g. Subsequently, the carbothermal reaction of CeO2 with C and oxygen produces CeC2, the reaction of Al with Ti produces Al3Ti and CeC2 reacts with dissolved Ti atoms to produce TiC. These reactions release a lot of heat [14,25]. Combined with literature reports [26,27], in the aluminum-rich melt, Al and Ti first form Al3Ti, and the first exothermic peak can be determined as the formation peak of Al3Ti. The formation temperatures are 783.2 and 791.7 ◦C, corresponding to the exothermic heat of 168.8 and 137.7 J/g. The second exothermic peak is the peak formed by TiC, the formation temperatures are 926.1 and 933.6 ◦C, and the corresponding exotherms of the reaction are 92.99 and 198.6 J/g. Compared with the Al-Ti-C system, it is found that the addition of CeO2 does not affect the melting point temperature of aluminum, but it reduces the latent heat of aluminum melting and accelerates the melting of aluminum. The smaller the size of CeO2, the lower the heat required for aluminum melting. Furthermore, the addition of CeO2 reduces the formation temperature and heat release of Al3Ti and TiC. The smaller the size of CeO2, the lower the formation temperature. The lower the heat released, the lower the temperature of the system, which affects the growth rate and growth habit of the second phase particles and is beneficial to obtaining Al3Ti and TiC with smaller size and regular shape.

**Figure 11.** DSC curves of Al-Ti-C-Ce with different CeO2 size of: 30 nm and 2–4 μm.

**Table 2.** The corresponding heat change of the reaction. ᶭᶭ ᶭ

#### *3.3. Refinement Performance Evaluation*

Figure 12 shows the macrographs of A 99.7 with different 0.3% master alloys. Figure 12a indicates that the unrefined pure aluminum is mainly composed of coarse equiaxed crystals in the center and columnar crystals in the middle and edges. From Figure 12b, it can be seen that Al-Ti-C master alloy has a significant refinement effect on A 99.7. After adding it into the melt of A 99.7, the coarse equiaxed crystals are refined into fine equiaxed crystals, while the range of columnar crystals is narrowed and the size is reduced. It can be seen from Figure 12c–f that the refining effect of Al-Ti-C-Ce master alloy is higher than Al-Ti-C master alloy, and the #2 alloy and #5 alloy have a better refining effect than #3 alloy and #4 alloy. Among them, the refining effect of #4 alloy with CeO2 size of 1 μm is the worst.

**Figure 12.** The macrographs of A 99.7 with 0.3% different master alloys: (**a**) unrefined; (**b**) #1; (**c**) #2; (**d**) #3; (**e**) #4; (**f**) #5.

Figure 13 shows the statistical results of the average grain size of the refined sample in Figure 12. It can be seen from Figure 13 that after 0.3% of Al-5Ti-0.62C, master alloys can refine the A 99.7 grain size from 1430 to 790 μm, while 0.3 wt.% #2 alloy refined the grain size of A 99.7 to 248 μm, and with the increased size of CeO2, grain size of A 99.7 tends to increase first and then decrease. The #5 alloy has a significant refinement effect on A 99.7, which can refine the grain size to 227 μm. The above analysis shows that the addition of CeO2 with a size of 2–4 μm to master alloy can promote its refining effect.

**Figure 13.** Average grain size of the refining samples in Figure 12.

#### **4. Conclusions**

**Author Contributions:** W.D. and T.C. conceived and designed the experiments; Y.M. and T.C. carried out the experiments and data collection; W.D., Y.M., T.C. and L.G. analyzed the data, Y.M. and W.D. contributed reagents/materials/analysis tools: Y.M., T.C. and L.G. wrote the paper. All authors have read and agreed to the published version of the manuscript.

**Funding:** National Natural Science Foundation of China (grant numbers 52161006 and 51661021) Industrial Support Plan Project of Gansu Provincial Department of Education (2021CYZC-23). China Postdoctoral Science Foundation, 2019M653896XB.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Exclude this statement.

**Acknowledgments:** This work was supported by the National Natural Science Foundation of China (grant numbers 52161006 and 51661021) and Industrial Support Plan Project of Gansu Provincial Department of Education (2021CYZC-23). The authors would like to acknowledge the China Postdoctoral Science Foundation, 2019M653896XB.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

## *Article* **Effect of Grain Refiner on Fracture Toughness of 7050 Ingot and Plate**

**Fang Yu 1,\*, Xiangjie Wang 1, Tongjian Huang 2 and Daiyi Chao 2**

**Abstract:** In this paper, two types of grain refining alloys, Al-3Ti-0.15C and Al-5Ti-0.2B, were used to cast two types of 7050 rolling ingots. The effect of Al-3Ti-0.15C and Al-5Ti-0.2B grain refiners on fracture toughness in different directions for 7050 ingots after heat treatment and 7050-T7651 plates was investigated using optical electron microscopy (OEM) and scanning electron microscopy (SEM). Mechanical properties testing included both tensile and plane strain fracture toughness (KIC). The grain size was measured from the surface to the center of the 7050 ingots with two different grain refiners. The fracture surface was analyzed by SEM and energy dispersive spectrometer (EDS). The experiments showed the grain size from edge to center was reduced in 7050 ingots with both the TiC and TiB refiners, and the grain size was larger for ingots with the Al-3Ti-0.15C grain refiner at the same position. The tensile properties of 7050 ingots after heat treatment with Al-3Ti-0.15C grain refiner were 1–2 MPa lower than the ingot with the Al-5Ti-0.2B grain refiner. For the 7050-T7651 100 mm thick plate with the Al-3Ti-0.15C grain refiner, for the L direction, the tensile properties were lower by about 10~15 MPa; for the plate with the Al-3Ti-0.15C refiner than plate with Al-5Ti-0.2B refiner, for the LT direction, the tensile properties were lower by about 13–18 MPa; and for the ST direction, they were lower by about 8–10 MPa compared to that of Al-5Ti-0.2B refiner. The fracture toughness of the 7050-T7651 plate produced using the Al-3Ti-0.15C ingot was approximately 2–6 MPa · √*m* higher than the plate produced from the Al-5Ti-0.2B ingot. Fractography of the failed fracture toughness specimens revealed that the path of crack propagation of the 7050 ingot after heat treatment produced from the Al-3Ti-0.15C grain refiner was more tortuous than in the ingot produced from the Al-5Ti-0.2B, which resulted in higher fracture toughness.

**Keywords:** Al-5Ti-0.2B; Al-3Ti-0.15C; 7050 ingot; 7050-T7651 plate; fracture toughness; grain size

#### **1. Introduction**

7050 Aluminum alloy have been widely used in the aerospace industry due to their low quench sensitivity, high fracture toughness, high fatigue crack growth rate (FCGR), and excellent stress corrosion resistance [1]. With the development of the aviation industry, the requirement for comprehensive performance of materials is getting higher, which means the raw material needs to have higher fracture toughness, higher FCGR, and high fatigue properties [2–4]. Normally, the tensile property levels decrease as the toughness level increases by alloying and heat treatment [5]. Therefore, the goal that most aluminum manufacturers want to pursue is to improve the fracture toughness of 7050 plates without the loss of the tensile property.

The effect of the microstructure on aluminum fracture toughness has been studied by many researchers. The main factors affecting fracture toughness [5,6] are grain size, recrystallization ratio, coarse insoluble phase, and strengthening precipitates. The relationship between recrystallization fraction (large angle grain boundary volume) and fracture toughness [7,8] was quantitatively analyzed. The higher the recrystallization fraction, the

**Citation:** Yu, F.; Wang, X.; Huang, T.; Chao, D. Effect of Grain Refiner on Fracture Toughness of 7050 Ingot and Plate. *Materials* **2021**, *14*, 6705. https://doi.org/10.3390/ma14216705

Academic Editor: Filippo Berto

Received: 11 September 2021 Accepted: 2 November 2021 Published: 7 November 2021

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**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

lower the fracture toughness. Hahn et al. [5] found that reducing the Fe, Si content and increasing the amount of smaller typical Cr, Mn, Zr-bearing particles could improve facture toughness, and the ratio of Zn:Mg [9] could also influence the fracture toughness. Zhang, Dumont et al. [10–12] investigated the effect of heat solution treatment process and quench parameters on the fracture toughness of 7050 and 7085 thick plate. Chen et al. [13] observed that the deformation texture can improve the fracture toughness and tensile properties of 7B50-T7751 plate effectively, and increase different directions of anisotropy. Other researchers also reported the effect of homogenization process [14] and aging parameter [15] on fracture toughness.

Although many experiments have been conducted to observe the factors that affect fracture toughness, there are still certain issues that have not been clarified. The grain size is known to be one of the major factors that affects fracture toughness [5], and grain refiners decide the final grain size and the quality of the ingot that is used for the manufacture of aviation products. A large number of studies have focused on the effect of grain refiners on the grain size and mechanical properties of ingots or products [16–21]. The effect of Al-5Ti-0.25C and Al-5Ti-0.2B additions on the grain size of different wrought alloys (AA1050, AA3004, AA5182, AA6063, AA7050, AA7475, AA8079) under different melting temperatures and addition rates was investigated [16] and it was found that the grain size decreased when increasing the rate of the refiner, and in all alloys, with the exception of AA3004, fine grain sizes (130 μm) were achieved with the addition of 1 kg/t. Nagaumi et al. [17] studied the effect of Al-3Ti-0.2B and Al-5Ti-1B on the microstructure of 7050 ingots and found that the refining fading of the Al-3Ti-0.2C refiner was more evident than that of Al-5Ti-1B grain refiner when the soaking time was increased. Huang et al. [18] noticed that compared with the 7050 alloy using the Al-5Ti-1B grain refiner, the distribution of the second phases present in 7050 aluminum alloy was more dispersed and uniform than the alloy using Al-5Ti-0.2C refiner. They also found that increasing the amount of Al-5Ti-0.2C grain refiner was more effective at easing the "Zr poisoning" phenomenon of 7050 alloy and could improve the strength and hardness while keeping good elongation. There are few reports [22] that discuss the relationship between grain refiner and fracture toughness of aviation aluminum products.

To examine the relationship between grain refiner and fracture toughness, a systematic study was conducted in 7050 ingots and 7050-T7651 100 mm thick plate with the grain refiner of Al-3Ti-0.15C and Al-5Ti-0.2B, respectively. The prepared plate was subjected to the same heat solution treatment and aging process. Microstructure characterization was then performed to evaluate the difference of the 7050 ingot and plate with two different grain refiners, and this study has guiding significance for industrial production of the 7050-T7651 plate, which needs higher fracture toughness.

#### **2. Materials and Methods**

The materials used for this study were commercial 7050 ingots with Al-3Ti-0.15C and Al-5Ti-0.2B grain refiner provided by Shan Dong Nanshan Aluminum Co., Ltd. at Longkou, China. The feeding amount for both Al-3Ti-0.15C and Al-5Ti-0.2B grain refiner was 2.5 kg/t. The ingot size was 440 mm × 1420 mm × 5000 mm and the composition is shown in Table 1.

**Table 1.** Chemical Composition of 7050 Aluminum Alloy (wt%).

After casting was completed, slices of the two different grain refined ingots were taken to evaluate the microstructure, grain size, and mechanical properties of the as-cast material according to the sampling diagram in Figure 1. The ingot samples were solution heat treated and artificially aged (heat solution treatment and age process were 479 ◦C × 50 min and 121 ◦C × 4 h + 163 ◦C × 20 h, respectively) prior to mechanical testing to eliminate the effects of the hot rolling process. The tensile properties and fracture toughness of ingot after heat treatment were then measured by an Instron 5985 tensile machine (INSTRON, Boston, MA, USA) and Instron 8802fatigue test machine (INSTRON, Boston, MA, USA), respectively. The fracture surface for ingots with different grain refiners were analyzed by scanning electron microscopy (SEM, FEI, Hillsboro, OR, USA).

**Figure 1.** Samples locations for 7050 ingot.

A 100 mm plate for the 7050 alloy was hot rolled from ingots that were cast with two different types of grain refining alloy. The first set of 7050 plates were produced using Al-3Ti-0.15C grain refined ingots, and the second set of 7050 plates were produced using Al-5Ti-0.2B grain refined ingots according to the preparation process shown in Figure 2. Ensuring the consistency of thermomechanical processing enables the impact of grain refining on the wrought microstructure and mechanical properties to be isolated. All of the TiC and TiB2 AA7050 ingots were homogenized together in the same furnace using a two-step process (step I: 465 ◦C soaking for 8 h, step II: 478 ◦C soaking for 20 h). After the homogenization process was finished, the ingots with Al-3Ti-0.15C and Al-5Ti-0.2B refiner were hot rolled to a 100 mm plate, then solution heat treated and artificially aged together to ensure consistent thermomechanical processes and to get the 7050 plate with T7651 temper (7050-T7651 plate).The heat solution heat treatment process was soaking 240 min at 479 ◦C, and the aging was a two-step process: Step I was soaking 4 h at 121 ◦C and Step II was holding 18 h at 163 ◦C.

**Figure 2.** Scheme of the preparation process of the studied 7050-T7651 plate.

Samples for the 100 mm 7050-T7651 plate were taken according to the requirement of the AMS2355 standard for tensile property and fracture toughness test, and the tensile properties and fracture toughness were tested on T/4 of plate by an Instron 5985 tensile machine and Instron 8802 fatigue test machine, respectively. Specimen orientations for fracture toughness are shown in Figure 3.

**Figure 3.** Specimen orientations used for fracture toughness test.

#### **3. Results**

#### *3.1. The Effect of Grain Refiner on 7050 Ingot Grain Size*

The grain size from the edge to the center of the 7050 ingot along with the casting direction was observed, five pictures were selected for each position, and the average grain size was analyzed by the intersection method according to ASTM E112 standard [23]. The microstructure and calculated grain size of the ingot with Al-5Ti-0.2B and Al-3Ti-0.15C grain refiner are shown in Figures 4 and 5, respectively. From Figure 4, the grain size from edge to center is increased, with a grain size of 95.99 μm, 101.04 μm, and 121.96 μm from edge to center for the ingot with Al-5Ti-0.2B grain refiner and 127.31 μm, 140.50 μm, and 165.32 μm with Al-3Ti-0.15C grain refiner, respectively. According to the study [17] by Hiromi, the average grain size of the 7050 ingot without grain refiner was 210 μm, and both the Al-5Ti-0.2B and Al-3Ti-0.15C grain refiner have refined the grain size.

**Figure 4.** Grain structure of 7050 ingot: (**a**) Al-5Ti-0.2B surface, (**b**) Al-5Ti-0.2B 1/4 thickness, (**c**) Al-5Ti-0.2B 1/2 thickness, (**d**) Al-3Ti-0.15C surface, (**e**) Al-3Ti-0.15C 1/4 thickness, and (**f**) Al-3Ti-0.15C 1/2 thickness.

**Figure 5.** Average grain size for ingot using different grain refiners.

#### *3.2. The Effect of Grain Refiner on Tensile Property of 7050 Ingot and 7050-T7651 100 mm Plate*

Table 2 shows the tensile properties of the 7050 ingot along with casting direction after heat treatment with Al-5Ti-0.2B and Al-3Ti-0.15C grain refiners. The tensile property of the 7050 ingot with Al-3Ti-0.15C grain refiner is a little lower than that of the ingot with Al-5Ti-0.2B grain refiner.

**Table 2.** Tensile Properties for 7050 Ingot after Treatment with Different Grain Refiners.

Table 3 depicts the tensile properties of 7050-T7651 100 mm thick plate with different grain refiners. According to the requirements of AMS 2355 standard [24], the mechanical properties at 1/4 thickness should be tested for plates with thickness of more than 38.1 mm. Therefore, for 7050-T7651 100 mm thick plate, the mechanical properties were tested at the location of 1/4 thickness to investigate the influence of different grain refiners on the mechanical properties. As can be seen from Table 3, compared with the addition of Al-5Ti-0.2B refiner, the tensile properties of 7050-T7651 100 mm thick plate with the addition of Al-3Ti-0.15C grain refiner in all directions are slightly reduced, and the elongation is not much different. For the L direction, the tensile property is lower by about 10–15 MPa for the plate with Al-3Ti-0.15C refiner than for the plate with Al-5Ti-0.2B refiner; for the LT direction, the tensile property is lower by about 13–18 MPa than that of Al-5Ti-0.2B refiner; and for the ST direction, the yield strength is lower by about 8–10 MPa than that of Al-5Ti-0.2B refiner.

**Table 3.** Tensile Properties for 7050-T7651 100 mm Plate with Different Grain Refiners.

#### *3.3. Effect of Different Grain Refiners on Fracture Toughness of 7050 Ingot and Plate*

Figure 6 shows fracture toughness test results for the 7050 ingot along with casting direction and 7050-T7651 100 mm plate at 1/4 thickness with different grain refiners. The fracture toughness of the ingot is increased by the addition of Al-3Ti-0.15C grain refiner. Comparing fracture toughness for different directions of the 7050-T7651 100 mm plate with Al-5Ti-0.2B grain refiner, plates with Al-3Ti-0.15C grain refiner have higher fracture toughness at L-T, T-L, and S-L directions, and the difference of the fracture toughness for the L-T direction for the 7050 product with different grain refiners is more obvious. The fracture toughness of the 7050-T7651 plate with Al-3Ti-0.15C grain refiner for the L-T direction is about 5 MPa · √*m* higher than for the plate with Al-5Ti-0.2B grain refiner.

**Figure 6.** Fracture toughness of 7050 ingot after heat treatment and 7050-T7651 100 mm plate for different grain refiners in different directions.

The fracture surfaces of the 7050 ingot fracture toughness specimens after heat treatment were analyzed by SEM to observe the effect of the grain refiners. Figure 7 shows fractography for different grain refiners of the ingots after heat treatment. The fracture surface exhibits different character for ingots with different grain refiners. Intergranular failure is observed linking shear planes from adjacent grains in ingots with Al-5Ti-0.2B grain refiner, whereas ingots with Al-3Ti-0.15C grain refiner exhibits transgranular failure via sheared planes covered with fine dimples. The crack path of the ingot with Al-3Ti-0.15C grain refiner is more tortuous.

The EBSD analysis results of different refiners for ingots after heat treatment at the crack tip are shown in Figure 8. As can be seen from Figure 8, the ingot containing Al-3Ti-0.15C refiner has intergranular and transgranular fracture, and the ingot produced by Al-5Ti-0.2B refiner presents intergranular fracture. Moreover, the ingot microstructure has a genetic effect: the grain size of the ingot with Al-3Ti-0.15C grain refiner is larger under the same refiner adding amount and casting process compared with Al-5Ti-0.2B grain refiner, which leads to increase in fracture toughness of the final product.

**Figure 7.** Fracture surfaces of ingot for different grain refiners: (**a**,**c**) Al-5Ti-0.2B grain refiner, (**b**,**d**) Al-3Ti-0.15C grain refiner.

**Figure 8.** EBSD analysis for 7050 ingot fracture surface at crack tip: (**a**,**b**) ingot with Al-3Ti-0.15C grain refiner, (**c**,**d**) ingot with Al-5Ti-0.2B grain refiner.

SEM analysis of the fracture surface for 7050-T7651 100 mm thick plate with different refiners were carried out. The samples in the T-L direction have both intergranular and transgranular fractures, and a large amount of second phases are distributed at the same time. The enlarged SEM photos of the fracture surfaces with different grain refiners are shown in Figure 9. The energy dispersive spectrometer (EDS) analysis shows both TiB2 and TiC particles have a granular shape, and the size of TiB2 particles is larger than that of TiC particles; the TiC particles were difficult to find under the scope of SEM. In addition, TiB2 particles are more aggregated than TiC particles, whereas TiC particles are more dispersed and have stronger resistance to crack propagation. The basic characteristics of fracture are basically consistent with the paper [25] researched by Zhang Xinming from Central South University.

**Figure 9.** T-L Fracture surface for 7050 plate: (**a**) Al-5Ti-0.2B grain refiner, (**b**) Al-3Ti-0.15C grain refiner (**c**,**d**) EDS result of particles from (**a**), (**e**,**f**) EDS result of particles from (**b**).

#### **4. Discussion**

The ingot is solidified from the surface to the center in turn. The surface layer has a high degree of supercooling, high nucleation rate, and relatively small grains because of the faster heat dissipation. Moreover, heat dissipation of the center is slow, the degree of undercooling is small, and the nucleation rate is low, resulting in larger grains in the

center. Literature [18] shows that the refining effect of Al-5Ti-0.2B is better than that of Al-3Ti-0.15C, primarily because of the microstructure of the refiner and the morphology and distribution of the refined particles. As shown in Figure 10, TiB2 crystal in the Al-5Ti-0.2B grain refiner has a closed hexagonal structure, which cannot provide a nucleation core for the matrix. The thin Al3Ti layer is enriched on the {0001} plane, which is used as the basis to promote α-Al nucleation. The nucleation directions of Al3Ti are <110>{112} and <210>{112}, respectively. During solidification, Al3Ti can nucleate in multiple directions and form a "halo" structure with "circle" shape; TiC crystal has face centered cubic (FCC) structure, which can be used as grain refining core for nucleation at the {111} plane and form "petalous" shape, and the number of nucleation is lower than that of the Al-Ti-B refiner. It is also related to the microstructure of the grain refiner and the morphology and distribution of the particles. The literature [17] shows that TiC particles in Al-3Ti-0.15C are unstable. When held in a soaking furnace for a long time, TiC particles will aggregate and react to form more stable Al4C3, whose refining effect is far less than that of TiC particles. Although for Al-Ti-B grain refiner, long soaking time also causes agglomeration and deposition of TiB2 particles, studies have shown [21] that Al-Ti-B is still effective in refining pure aluminum, even if the holding time in the furnace is as long as 24 h. Combined with the above factors, the grain refinement effect of Al-3Ti-0.15C is not as good as that of Al-5Ti-0.2B. Therefore, with the same amount of addition, the grain size of the ingot with Al-5Ti-0.2B grain refiner is smaller.

**Figure 10.** Grain refiner nucleation schematic diagram: (**a**) Orientation relationship between Al, Al3Ti, and TiB2; (**b**) Nucleation diagram of Al-5Ti-0.2B and Al-3Ti-0.15C refiner during solidification.

If the grain size is small enough, plasticity will be enhanced without detrimental, low energy, intergranular fracture for thin products under plane stress. However, for thick products under plane strain, fracture is usually controlled by coarse particles, and a recrystallized grain structure is preferable. Teleshov [26] studied the effect of grain size on fracture toughness and strength of AK4-1CH extrusion products. It was shown that the larger grain size products have higher fracture toughness than the smaller grain size products in L-T, T-L, and S-L directions, especially in the T-L direction. By comparing the microstructure, properties, and the distance between the vertical crack plane and the second phase in the propagation direction, it was found that the main reason for the different fracture toughness of AK4-1CH products may be the smaller extent of grain boundary in the path of plastic deformation. To overcome this deformation, the deformation in neighboring grains is needed. However, the stress generated far from the crack tip is not enough, which leads to the increase in the size of the plastic deformation zone, resulting in the increase in KIC. According to the study [27] of Krasovskii, the depth of the plastic deformation zone is related to the fracture toughness and yield strength of the product. According to the equation λ = 0.128 ∗ (KIC/δ0.2) 2, the value for the ingot with different grain refiners is 1.080 mm and 0.787 mm, respectively, which are larger than the average grain size. Fracture path for the 7050 ingot with different grain refiners is shown in Figure 11. T range of the

plastic deformation zone with large grain size was large, which affected the fracture mode at the crack tip.

**Figure 11.** Fracture path for 7050 ingot with different grain refiners: (**a**) Al-3Ti-0.15C (**b**) Al-5Ti-0.2B.

As discussed in Section 2, the recrystallization ratio is one of the factors affecting fracture toughness: the higher the recrystallization fraction, the lower the fracture toughness. The recrystallization ratio of the 7050-T7651 plate in different thickness positions with Al-5Ti-0.2B and Al-3Ti-0.15C grain refiners are shown in Figure 12. From surface to 1/2 thickness, the recrystallization ratio for the 7050 plate with Al-5Ti-0.2B grain refiner is 11.2%, 10.3%, and 2.4%, respectively, and 12.3%, 8.1%, and 1.5% with Al-5Ti-0.2B grain refiner. There are no obvious differences in the recrystallization ratio for the 7050-T7651 plate with different grain refiners.

**Figure 12.** Recrystallization ratio of 7050-T7651 plate with different grain refiners.

The Fe and Si bearing phase formed during solidification is brittle and not coherent with the matrix, easily forms pores, and becomes a source of crack propagation, thus strongly reducing the fracture toughness of the alloy. The size of the precipitate and the interface binding force between the precipitate and the matrix have a great influence on the fracture toughness. The larger the size of the precipitate, the weaker the binding force between the phase and the matrix, which leads to the crack occurrence and the inferior fracture toughness. Gokhale, Deshpande et al. [7,28] quantitatively analyzed the relationship between microstructure and fracture toughness, and found that the larger the surface area of the second phase particle per unit volume, the lower the fracture toughness. Therefore, in order to improve the fracture toughness of 7050 products, the content of Fe and Si and the size of the dispersion phase must be strictly controlled. In the 7050-T7651 plate with Al-5Ti-0.2B grain refiner, TiB2 particles are easy to gather in the process of casting and have a genetic effect; TiB2 particles act as heterogeneous nucleation to precipitate in the aging process. The large precipitate phase will be become the source of cracks, which may reduce the local plastic deformation ability and also the fracture toughness.

#### **5. Conclusions**

A study has been performed on the effect of grain refiner on the tensile property and fracture toughness of 7050 ingots and 7050-T7651 100 mm thick plate. The following conclusions may be drawn from this work.

The experiments showed that the grain size from surface to center of the ingot with the addition of Al-5Ti-0.2B grain refiner was 95.99 μm,101.04 μm, and 121.96 μm, respectively, and 127.31 μm, 140.50 μm, and 165.32 μm with the addition of Al-3Ti-0.15C grain refiner. The refining ability of Al-5Ti-0.2B grain refiner was better than that of Al-3Ti-0.15C.

For both 7050 ingots and 7050-T7651 100 mm thick plate, the tensile property with Al-5Ti-0.2B grain refiner was higher than that of Al-3Ti-0.15C grain refiner. However, the fracture toughness of 7050 ingots and 7050-T7651 plates made with Al-3Ti-0.15C grain refiner was higher than that made with Al-5Ti-0.2B grain refiner, and fracture path for the 7050 ingot with Al-3Ti-0.15C after heat treatment was more tortuous than for the ingot with Al-5Ti-0.2B.

The fractured surface for the 7050-T7651 plate showed both TiB2 and TiC particles have a granular shape, and the size of the TiB2 particles was larger than that of the TiC particles. TiC particles were hard to find under the scope of SEM, whereas TiB2 particles were easy to find and easy to aggregate. The combination of grain size and uniform distribution of the second phase results in higher fracture toughness for 7050 ingots and 7050-T7651 plate with Al-3Ti-0.15C grain refiner.

**Author Contributions:** Conceptualization, F.Y. and X.W.; methodology, F.Y. and T.H.; software, F.Y.; validation, X.W., D.C. and T.H.; formal analysis, F.Y.; investigation, F.Y.; resources, X.W.; data curation, D.C.; writing—F.Y.; writing—review and editing, F.Y., D.C. and X.W. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Fundamental Research Funds for the National Natural Science Foundation of China (U1708251); the Key Research and Development Program of Liaoning (2020JH2/10700003).

**Institutional Review Board Statement:** The study was conducted according to the guidelines of the Northeastern University, and approved by the School of Material Science and Engineering.

**Informed Consent Statement:** Informed Consent was obtained from all subjects involved in the study.

**Data Availability Statement:** The data presented in this study are available upon request.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

## *Article* **In Situ Observation of the Tensile Deformation and Fracture Behavior of Ti–5Al–5Mo–5V–1Cr–1Fe Alloy with Different Microstructures**

**Suping Pan 1,2, Mingzhu Fu 1, Huiqun Liu 1,3,\*, Yuqiang Chen 4,\* and Danqing Yi 1,3**

**Abstract:** The plastic deformation processes and fracture behavior of a Ti–5Al–5Mo–5V–1Cr–1Fe alloy with bimodal and lamellar microstructures were studied by room-temperature tensile tests with in situ scanning electron microscopy (SEM) observations. The results indicate that a bimodal microstructure has a lower strength but higher ductility than a lamellar microstructure. For the bimodal microstructure, parallel, deep slip bands (SBs) are first noticed in the primary α (αp) phase lying at an angle of about 45◦ to the direction of the applied tension, while they are first observed in the coarse lath α (αL) phase or its interface at grain boundaries (GBs) for the lamellar microstructure. The β matrix undergoes larger plastic deformation than the αL phase in the bimodal microstructure before fracture. Microcracks are prone to nucleate at the αp/β interface and interconnect, finally causing the fracture of the bimodal microstructure. The plastic deformation is mainly restricted to within the coarse αL phase at GBs, which promotes the formation of microcracks and the intergranular fracture of the lamellar microstructure.

**Keywords:** Ti–5Al–5Mo–5V–1Cr–1Fe alloy; in situ observation; slip band; microcrack; fracture mechanism

#### **1. Introduction**

Due to their strength, corrosion resistance, and heat resistance, near-β-titanium alloys have been widely used in aerospace and automotive engineering as structural materials [1]. With the continuous progress of aerospace and automotive technology, more challenging requirements are imposed on the mechanical properties of titanium alloys [2]. The design of titanium alloys with better mechanical properties is therefore the focus of many studies [3–6].

Previous studies found that the mechanical properties of titanium alloys are closely related to the morphology and distribution of the α phase [7–9]. Two types of typical microstructure (bimodal and lamellar microstructures) can be obtained in titanium alloys by different heat treatments [10]. Many studies concerning the microstructure–property relationship of titanium alloys showed that bimodal variants have a higher ductility and lower strength than their lamellar counterparts [11–13]. Qin et al. [14] found that the yield strength (σ0.2) of lamellar variants reaches 1900 MPa while that of bimodal alloys is around 1600 MPa. Zheng et al. [15] found that the elongation of bimodal variants is 11.5–14.5% and that of lamellar alloys is 4.5–9.5%. Therefore, the microstructure plays an essential role in determining the mechanical properties of titanium alloys.

Many studies have been performed to reveal the essential mechanism underpinning the different mechanical properties of titanium alloys induced by bimodal and lamellar

**Citation:** Pan, S.; Fu, M.; Liu, H.; Chen, Y.; Yi, D. In Situ Observation of the Tensile Deformation and Fracture Behavior of Ti–5Al–5Mo–5V–1Cr–1Fe Alloy with Different Microstructures. *Materials* **2021**, *14*, 5794. https:// doi.org/10.3390/ma14195794

Academic Editors: Andrey Belyakov and Carmine Maletta

Received: 21 July 2021 Accepted: 29 September 2021 Published: 3 October 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

microstructures. Wu et al. [16] found that the higher ductility of bimodal alloys can be mainly attributed to the more numerous deformation mechanisms available to the equiaxed primary α phase (αp) in bimodal alloys (including dislocation slips, twins, and shear bands) than those of the α lath (αL) in lamellar structures. Some believed that the tensile fracture of near-β-titanium alloys is sensitive to micro-void nucleation. Qin et al. [17] found that the tensile fracture of lamellar alloys is a process of the initiation of nano-scale voids, followed by their growth and coalescence in the deformation band, and transgranular shearing. The reason for the formation of micro-voids is the stress concentration caused by the difference in the strength of the α phase and β matrix. Stress concentration at the GBs is derived from the precipitation of the αL phase along the α/β interface, which results in micro-void nucleation. This behavior is mainly responsible for the low ductility of lamellar alloys. Prior research into the deformation and properties of materials is based on ex situ test methods, which make it difficult to provide direct evidence of the deformation and fracture behavior; therefore, detailed information concerning the evolution of a microstructure, and its corresponding effect on the fracture process under tension, remains unclear.

Recently, in situ tensile SEM observation has become a powerful and effective tool to evaluate the deformation behavior from a microstructural perspective by capturing the microstructural evolution dynamically [18,19]. Huang et al. [20] indicated that localized stress concentration at the GBs derived from the geometric incompatibility between neighboring α grains was mainly responsible for microcrack formation in a Ti–6Al alloy. Zhang et al. [21] studied the deformation mechanism of a Ti–5Al–2Sn–2Zr–4Cr–4Mo alloy with a bimodal microstructure, implying that the αp had high compatibility of deformation and the slip line in the αp phase was the primary deformation mechanism. Shao et al. [22] found that microcracks were primarily initiated along the αL phase at the edges of the sample. To date, however, the essential mechanism of the effect of microstructure on the mechanical properties of titanium alloys remains unclear, which hinders any further attempts to improve the mechanical properties of alloys.

In this study, Ti–5Al–5Mo–5V–1Cr–1Fe alloy samples with bimodal and lamellar microstructures were prepared through heat treatments. Their room-temperature tensile deformation process and fracture behavior were monitored in real time by in situ SEM observations. On this basis, the aim of this work is to ascertain the essential effects of microstructure on the mechanism of tensile deformation and fracture of titanium alloys. The results of this work are expected to provide a basis for future improvement in the mechanical properties of such titanium alloys.

#### **2. Materials and Methods**

#### *2.1. Materials*

The as-received material in this study was a forged Ti–5Al–5Mo–5V–1Cr–1Fe alloy provided by Baoti Group Ltd (Baoji, China) as cuboidal specimens (98.0 mm in length, 20.0 mm in width, and 7.0 mm in height) with a nominal composition of 5.07% Al, 4.81% Mo, 4.74% V, 1.06% Fe, 0.95% Cr, and the rest Ti (all in wt%). The α + β/β transition temperature (*Tα + β*→*β*) of this alloy is about 865 ◦C.

To obtain the desired microstructure, two as-received cuboids were, respectively, solution-treated at 830 ◦C (below *Tα + β*→*β*) and 895 ◦C (above *Tα + β*→*β*) for 2 h. Then, these cuboids were cooled in a furnace (FC) to 750 ◦C and held for 2 h before air cooling (AC). Thereafter, they were aged at 600 ◦C for 8 h, before air cooling (AC) to room temperature (Figure 1).

**Figure 1.** Heat treatment routes for bimodal and lamellar samples.

#### *2.2. In Situ Tensile Test*

The sample for the in situ tensile test was cut by wire-electrode cutting with a gauge length of 1.5 mm, a gauge width of 1.5 mm, and a thickness of 1.0 mm (Figure 2). It was mechanically polished using emery papers with SiC (5 μm, 3 μm, 1 μm, and 0.25 μm). Then, it was chemo-mechanically polished using Al2O3 (0.04 μm) suspension to remove the workhardening resulting from previous mechanical polishing. A micro-stage (Figure 2a) and Mini MTS (Liweiauto Ltd., Hangzhou, China) controller (with a maximum load capacity of 2500 N) were employed to clamp the sample and control the tensile strain rate, respectively. The Mini MTS system was adopted to measure the mechanical properties of in situ samples. Before in situ testing, the sample was mechanically polished and then stretched at a speed of 1.5 μm/s at room temperature in vacuum.

**Figure 2.** (**a**) Mechanical test bench; (**b**) geometry and dimensions of in situ tensile test specimens at room temperature (units: mm).

Processes during the in situ tensile deformation in bimodal and lamellar microstructures were observed using a JSM-5600 field-emission gun SEM (JEOL Ltd., Tokyo, Japan) with an accelerating voltage of 15 kV. Several interruptions were allowed by the loading system during tensile testing: this allowed the load to be held while capturing the SEM micrographs, after which the tensile test was resumed from the same applied load and displacement at the same rate. Three repeated in situ tensile tests were carried out under each set of experimental conditions, and typical results were provided. The tensile direction for the corresponding SEM images printed herein was parallel to the vertical direction.

Electron backscattered diffraction (EBSD) measurements were conducted before and after in situ tensile tests using an AZtec system (Oxford Instruments Group, Oxford, UK) coupled to a Hitachi-Regulus 8230 cold field emission SEM (Hitachi High-Technologies Corporation, Tokyo, Japan). The operating voltage used was 20 kV to optimize the quality

of the diffraction patterns. The EBSD samples were electropolished using a solution of 8% perchloric acid (HClO4) and 92% CH4O at −25 ◦C.

#### **3. Results**

#### *3.1. Microstructure of As-Heat-Treated Samples*

Figure 3a shows backscattered electron (BSE)-SEM images of bimodal samples before in situ tensile testing. Since this sample was solution-treated at a temperature below *Tα + β*→*β*, some coarse globular αp phases were maintained, which were evenly distributed in the β matrix. For the lamellar microstructure seen in Figure 3b, αp was completely dissolved during solution treatment at a temperature above *Tα + β*→*β* (i.e., 895 ◦C for 2 h), while coarse αL, formed during subsequent low-temperature aging treatment, was distributed along GBs of the β matrix.

**Figure 3.** BSE-SEM images of samples before stretching in situ; (**a**) bimodal microstructure; (**b**) lamellar microstructure.

#### *3.2. The Stress–Displacement Curves during In Situ Testing*

The stress–displacement curves of bimodal and lamellar microstructures under in situ stretching are demonstrated in Figure 4. The drops in the curves are caused by slight stress relaxation during the pauses for SEM imaging, in which three typical drops for each sample are marked, respectively (A, B, and C for the bimodal microstructure and A', B', and C' for the lamellar microstructure).

**Figure 4.** The stress–displacement curves of samples with bimodal and lamellar microstructures stretched in situ at room temperature.

As can be seen from Figure 4, σ0.2 and the ultimate tensile stress (σth) of the bimodal microstructure are 1041.8 and 1220.5 MPa, which are much lower than those of specimens with a lamellar microstructure (σ0.2 = 1102.7 MPa, σth = 1441.3 MPa). The maximum tensile displacement of the bimodal microstructure is about 872 μm, while that of the lamellar microstructure is about 736 μm. This result indicates that the lamellar microstructure is stronger, but less ductile than the bimodal microstructure, which is consistent with findings from previous studies [11–15].

#### *3.3. Microstructure Evolution during In Situ Stretching*

#### 3.3.1. Microstructure Evolution of Bimodal Microstructure

Figure 5a illustrates the in situ SEM images of the bimodal microstructure at position A. At this stage, some parallel and deep SBs can be seen in a small number of αp phase regions. These SBs generally lie at an angle of 41◦ to 49◦ with the tensile direction. As shown in Figure 5c, most of them are quite short (only several microns) and strictly confined within a single αp phase, while some of them not only cross the whole αp phase grain, but also pass through the αp/β interface and enter the region containing the β matrix (Figure 5b). In addition, there is a certain region of distortion arising at the αp/β interface (Figure 5d), which may be attributed to the deformation incompatibility between αp and β due to their different crystal structures.

**Figure 5.** (**a**) In situ SEM images of the bimodal microstructure at position A; (**b**) magnified image showing SBs passing through the αp/β interface; (**c**) SBs formed in the region containing the αp phase; (**d**) distortion at the αp/β interface region.

Figure 6 illustrates BSE-SEM images of bimodal samples stretched at position B. Obvious necking appears (Figure 6a) and some microcracks form on the edge of the sample (Figure 6b) caused by the increasing strain. The microcrack tends to propagate along the αp/β interface and gradually grows to the center of the sample, as shown in Figure 6b. Noticeable SBs can be found in the β matrix adjacent to regions of αp phase and the distortion at the αp/β interface increases in severity (Figure 6d). This indicates that a higher stress concentration arises at this region near the αp/β interface. In addition,

Figure 6c shows that the SBs in the αp phase become deepened and some tiny microcracks are also initiated in the β matrix close to regions containing the αp phase.

**Figure 6.** (**a**) In situ SEM images of the bimodal microstructure at position B; (**b**) magnified image showing microcracks formed on the edge of the sample; (**c**) microcracks initiated from the αp/β interface; (**d**) distortion was aggravated around the αp/β interface.

Figure 7 illustrates in situ SEM images of the bimodal microstructure at position C (after fracture). The fracture surface is relatively rough (Figure 7e) and significant necking can be observed at the region close to the fracture surface (Figure 7c,f). As shown in Figure 7b,d, many microcracks form in the region close to the fracture surface: these nucleate at and generally propagate along the αp/β interface (the crack path is tortuous). Although the number of SBs increases, the αp phase still maintains a granular shape (Figure 7d) and grains are wrapped by the significantly distorted β matrix (Figure 7d). As can be seen in Figure 7a, the morphology of the β matrix near the fractured zone becomes streamlined in shape, implying that it undergoes significant plastic deformation before fracture. Therefore, the β matrix undergoes greater deformation than the αp phase during in situ stretching; in addition, this deformation is relatively uniformly distributed in the β matrix, probably due to the excellent deformation compatibility.

To assess the deformation behavior of the bimodal microstructure during in situ tensile loading, an area of the sample was selected for tracking, and the SEM images thereof at different stages were recorded. Figure 8a,c show the SEM images of the same area on a specimen with a bimodal microstructure surface at the stage of positions A and B, respectively: with increasing strain, the sample surface became significantly rougher and took on an undulate appearance (Figure 8c). To determine the strain in this local area, the changes in distance between two αp phase grains on the sample surface were measured. Stretched from position A to position B, the distance increased from 135.71 μm (position A) to 167.85 μm (position B), indicating some 23.68% of the plastic deformation appears along the tensile direction (*εL*) in this region.

**Figure 7.** (**a**) The streamlined shape of the β matrix near the fractured zone at position C; (**b**) a crack propagated along the αp/β interface; (**c**,**f**) SEM images showing the rough fracture surface of the sample; (**d**) magnified image showing intensified SB formation in the αp phase and microcracking at the αp/β interface; (**e**) macroscopical SEM images of the bimodal microstructure at position C.

Figure 8b,d are magnified images of Figure 8a,c, respectively. Although the sample underwent severe plastic deformation, the distribution of SBs in the β matrix remained relatively uniform. For further analysis, 25 αp phase grains were selected for calibration (and assigned serial numbers 1 to 25). The size changes of these 25 αp phase grains along the tensile direction from positions A to B were statistically studied (Table 1); most αp phase regions have *εL* values of less than 10%, and their average *εL* value is 9.72%, which is only 41% of the average *εL* value over the region. This result demonstrates that during tensile deformation, the β matrix was subject to larger plastic deformation than the αp phase.

**Figure 8.** In situ SEM images of bimodal samples; (**a**) a selected area at position A; (**b**) magnified image showing the morphologies of αp phase at position A; (**c**) the area in (**a**) stretched to position B; (**d**) magnified image illustrating the deformation within αp phase at position B.

**Table 1.** Calculation of *εL* values for αp phase (Figure 8) from positions A to B.

#### 3.3.2. Evolution of the Lamellar Microstructure

Figure 9 shows BSE-SEM images of specimens with a lamellar microstructure during in situ stretching at position A' on the stress–displacement curve (Figure 4). As shown in Figure 9d, some parallel SBs form within or at the boundaries of the αL phase at an angle of about 45◦ to the tensile direction. Differing from those in specimens with a bimodal microstructure, in which such SBs measure only several microns, these SBs aligned along the length of the αL phase are significantly longer and can always grow to several tens of microns.

**Figure 9.** (**a**) In situ SEM images of lamellar microstructure at position A'; (**b**,**c**) magnified image showing several SBs in an αL phase or at its interface; (**d**) magnified image showing plentiful SBs in a long coarse αL phase at a GB at an angle of about 45◦ to the tensile direction.

Figure 10 illustrates BSE-SEM images of lamellar samples during in situ stretching at position B'. As the strain increases, SBs gradually extend and connect with each other along the length of αL at which some extremely long SBs are formed. As shown in Figure 10d, the connected SBs in the αL at the grain boundary are over one hundred microns in length, besides which certain microcracks evolving from SBs can also be found in the αL grains or at their interfaces.

**Figure 10.** (**a**) In situ SEM images of lamellar microstructure at position B'; (**b**,**c**) magnified image showing several microcracks formed in αL phase or at their interfaces; (**d**) magnified image showing SBs connected with each other and forming an SB with a length of 127.51 μm.

> Figure 11 exhibits in situ SEM images of lamellar samples during stretching at position C'. No obvious necking occurs in the sample until it fractures, which indicates that the lamellar microstructure undergoes less plastic deformation than the bimodal microstructure. Although there are slight SBs in the β matrix adjacent to the fracture surface, they are fewer in number than in specimens with a bimodal microstructure. Judging from Figure 11c,d, specimens with a lamellar microstructure generally fracture along the GBs, which results in a sharp fracture surface.

**Figure 11.** (**a**) In situ SEM images of the lamellar microstructure at position C'; (**b**,**d**) magnified images showing many small, shallow dimples on the fracture surface; (**c**) magnified image showing the sample with a shape fracture surface generally breaking along the GB; (**e**) some short crossed SBs are shown in the β matrix adjacent to the fracture surface.

> As shown in Figure 12a,c, the changes in a selected area of the lamellar microstructure from position A' to position B' are tracked. The changes in roughness of the sample surface are small and there are fewer SBs compared with those in specimens with a bimodal microstructure (Figure 11). The εL value in this local area from position A' to position B' is found to be 1.33% (less than that in specimens with a bimodal microstructure).

**Figure 12.** In situ SEM images of the lamellar microstructure; (**a**) a selected area at position A'; (**b**) magnified image showing the morphologies of αL phase at position A'; (**c**) the area in (**a**) stretched to position B'; (**d**) magnified image illustrating the deformation within αL phase at position B'.

Figure 12b,d demonstrate magnified versions of Figure 12a,c, respectively. To quantify the deformation behavior, 12 αL phase grains were selected for calibration and assigned serial numbers 1 to 12. As shown in Figure 12b, SBs are mainly found in the large αL phase (αL12), which can extend along the GB, while they are scarcely found in the β matrix or small αL phase regions. With the increase in strain from position A' to position B', SBs in αL12 deepen, while the microstructure in the β matrix and small αL phase changes little (Figure 12d). The εL values of the selected 12 αL phase grains deformed from position A' to position B' were statistically analyzed (Table 2). According to the data in Table 2, the relatively small αL phase (αL1-αL11) has quite small εL values (less than 1.0%), while the εL value of the long αL phase at the GB (αL12) is found to be 5.79%, which is 3.35 times greater than the εL value across this area; the deformation in the lamellar microstructure is mainly concentrated in the large αL phase found at the GB (i.e., αL12).

**Table 2.** Calculation of *εL* values of αL phase (Figure 12) from positions A' to B'.

**Table 2.** *Cont.*

#### **4. Discussion**

Recently, investigators have found that mechanical behavior and its related mechanism of action are sensitive to the initial microstructure of titanium alloys. Specimens with a bimodal microstructure are found to have a lower strength but higher ductility than those with a lamellar microstructure [11–15], while the essential reason for the differences in mechanical behavior between bimodal and lamellar microstructures is still hotly debated. Huang et al. [23] found that the αp phase in specimens with a bimodal microstructure plays a major role in accommodating the plastic strain of titanium alloys due to its good compatibility during deformation, while Tan et al. [24] stated that cracks are readily initiated at SBs in the αp phase, which lies at the crux of the tensile deformation. Liu et al. [25,26] considered that a high stress concentration at the αL/β grain boundary results in intergranular fracture and low ductility of specimens with a lamellar microstructure, whereas Qin et al. [27] found that the crack nucleates inside the β grains and will spread under high tensile stress without hindrance in the larger β grains, leading to the low plasticity of specimens with a lamellar microstructure. In this study, their quite different mechanical properties were found to be essentially attributed to different microstructural evolutions during tensile loading.

#### *4.1. Deformation Mechanisms and Microstructural Evolution of the Bimodal Microstructure*

Based on the results of in situ SEM observation, the deformation mechanisms and microstructural evolution of the bimodal microstructure are shown schematically in Figure 13.

The bimodal microstructure contains coarse globular αp grains distributed in the β matrix (Figure 13a). Under a relatively small strain (Figure 13b), many parallel, deep SBs are formed inside some αp phase regions due to the limited slip systems of αp and the stress concentration caused by their relatively large size. As previously proved by Semiatin et al. [28], the ratio of critical resolved shear stress in the α phase at room temperature was 1:0.7:3.2 for basal ({0001}<11−20>), prismatic ({1−100}<11−20>), and pyramidal slip ({1−101}<11−20>), respectively. This indicates that pyramidal slip is difficult to take place and the αp phase is more likely to slip along its basal or prismatic plane at room temperature. Figure 14 demonstrates four examples of SB identification for a bimodal microstructure after in situ stretching. These SBs essentially correspond to prismatic or basal slip systems and to single slip behavior with a relatively large Schmid factor (SF) (SF > 3.7), which agrees well with the results of Semiatin et al. [28].

**Figure 13.** Schematic mechanism showing the evolution of bimodal microstructure during in situ stretching; (**a**) the initial microstructure; (**b**) at a relatively low strain; (**c**) at a relatively high strain; (**d**) schematic mechanism of deformation of bimodal samples after fracture.

**Figure 14.** Examples of identification of activated slip systems on the αp phase in bimodal microstructure; (**a**,**b**) a basal slip activates (SF = 0.38) in the αp phase; (**c**,**d**) a basal slip activates (SF = 0.37) in the αp phase; (**e**,**f**) a prismatic slip activates (SF = 0.45) in the αp phase; (**g**,**h**) a prismatic slip activates (SF = 0.46) in the αp phase.

In addition, as shown in Figure 15a, the orientations of αp phases are randomly distributed in the bimodal microstructure. For the same αp phase, there is a large difference in SF values for basal and prismatic slips (Figure 15b,c); the one with the maximum SF is supposed to be the easiest to activate [21]. Therefore, it is difficult for the αp phase to activate both basal and prismatic slips simultaneously. SBs are prone to occur along a single basal or prismatic plane of the αp phase, which is oriented solely along the maximum shear stress direction, i.e., at an angle of 45◦ to the applied tension. Differing from the αp

phase, there are significantly more slip systems in the β phase, among which {1−10}<111>, {11−2}<111>, and {12−3}<111> slips were thought to be the three typical cases [28,29]. As presented in Figure 15d,f, there is no significant difference in the SF values for these three types of slips in the β phase. Thus, it is difficult to observe β phase regions' slip strictly along one single plane especially under a small strain.

**Figure 15.** EBSD images of bimodal microstructure before stretching in situ; (**a**) inverse pole-figure (IPF) map; (**b**) SF map of α for basal slip; (**c**) SF map of α for prismatic slip; (**d**) SF map of β for {1−10}<111> slip; (**e**) SF map of β for {11−2}<111> slip; (**f**) SF map of β for {12−3}<111> slip.

In the bimodal microstructure, although a small number of SBs in the αp phase can pass through the αp/β interface, most are restricted at the interface between β and αp and confined to within a single αp phase grain. This will lead to an increased stress concentration at the αp/β interface, therefore distorting the interface. As the strain increases (Figure 13c), the αp phase is elongated slightly along the tensile direction (Figure 8).

β is softer than αp due to lower concentrations of solute Al in β [30], thus β should bear more significant deformation than the αp phase at the same stress, while due to the more numerous slip systems and greater deformation coordination in BCC β, parallel, deep SBs may be less likely to form in the β phase when the strain is relatively small (Figures 6 and 7). Caused by the different deformation behaviors between αp and β, stress concentration at the interface gradually increases with increasing strain, finally generating microcracks. Then, with the further increase in strain (Figure 13d), the number of SBs in the αp phase increases slightly due to its more limited slip system, while that in the β matrix increases to a much greater extent. This further aggravates the stress concentration at the αp/β interface, making the microcrack propagate along the αp/β interface and into the β matrix. As the microcracks grow, they gradually come closer together, whereupon they tend to interconnect to form a main crack. Additionally, as the main crack grows, it is likely to bridge these microcracks formed at the αp/β interfaces. This finally leads to the zig-zag crack path and rough fracture surface of the sample since the αp grains are randomly distributed in the β matrix (Figure 13d).

It should also be pointed out that, because of the connected distribution of the β phase and isolated distribution of the αp phase in specimens with a bimodal microstructure, localized deformation readily propagates into the surrounding area by way of the soft β phase, with only a minor role played by the harder αp phase. This ensures relatively uniform deformation, giving rise to the excellent ductility of the bimodal microstructure.

#### *4.2. Deformation Mechanisms and Microstructural Evolution of the Lamellar Microstructure*

Figure 16 schematically presents the deformation mechanisms and microstructural evolution of the lamellar microstructure based on the results of in situ SEM observation.

**Figure 16.** Schematic mechanism showing the evolution of lamellar samples during in situ stretching; (**a**) initial microstructure; (**b**) at a relatively low tensile displacement; (**c**) at a relatively high tensile displacement; (**d**) schematic mechanism of the deformation of lamellar samples after fracture.

Differing from the bimodal microstructure, many long and coarse lamellar αL grains, rather than globular αp regions, are seen in the lamellar microstructure: these are mainly distributed at or near GBs (Figure 16a). Since these αL phase regions are much harder than the β phase [31], they can exert a strong fencing effect and therefore separate each β grain into a relatively isolated region between which dislocation cannot easily pass. This in turn leads to a great stress concentration at the αL phase. Therefore, deep SBs are first observed in the coarse αL phase at GBs under a relatively low tensile strain (Figure 16b). As shown in Figure 17, these SBs correspond to prismatic slip behavior (SF > 4.2) and generally lie along the plane at an angle of approximately 45◦ to the direction of the applied tension (i.e., the direction of maximum shear stress).

Meanwhile, the deformation in the β matrix is much smaller due to the strong fencing effect of the αL (Figure 9). With increasing strain, the number of SBs in the large αL phase increases while that in the β matrix remains low (Figure 10), suggesting that the deformation in lamellar microstructures is non-uniform and mainly concentrated in the large αL grains at GBs. As a result of this inhomogeneous deformation, microcracks are readily initiated from, and propagate along, the SBs (Figure 16c). The formation of microcracks in these αL regions in turn produces a greater stress concentration, releasing stress accumulation in the β matrix. This aggravates the inhomogeneity of the deformation of specimens with a lamellar microstructure. Finally, the sample fractures along the large αL grains at GBs, leading to the low plasticity of the lamellar microstructure (Figure 16d).

**Figure 17.** Examples of identification of activated slip systems on αL phase in lamellar microstructure; (**a**,**b**) prismatic slips activate in two αL phase with SF values of 0.44 and 0.42, respectively; (**c**,**d**) prismatic slips activate in two αL phases with SF values of 0.43 and 0.45, respectively.

#### **5. Conclusions**

In this study, the microstructural evolution and fracture mechanisms of a Ti–5Al–5Mo– 5V–1Cr–1Fe alloy with bimodal and lamellar microstructures were investigated through in situ tensile SEM observation. The following conclusions can be obtained:

**Author Contributions:** Investigation, S.P. and Y.C.; Data acquisition Y.C., S.P. and M.F.; Formal analysis, M.F., H.L., D.Y. and S.P.; Resources, H.L. and Y.C.; Writing—original draft, S.P. and M.F.; Conceptualization, D.Y. H.L. and Y.C. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the State Key Laboratory of Powder Metallurgy (Grant No. 10500-621022001), Central South University, Changsha, China.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Informed consent was obtained from all subjects involved in the study.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

## *Article* **Cobalt Content Effect on the Magnetic Properties of Ni50***−***xCoxMn35.5In14.5 Annealed Ribbons**

**Łukasz Dubiel 1,\*, Andrzej Wal 2, Ireneusz Stefaniuk 3, Antoni Zywczak ˙ 4, Piotr Potera 2 and Wojciech Maziarz 5**

**Abstract:** We present a study of the annealing effect and its influence on magnetic and structural properties for a series of Heusler alloys Ni50−xCoxMn35.5In14.5 (x = 0, 3, 5) prepared in ribbon form. We studied the morphology and composition using scanning electron microscopy (SEM) equipped with an X-ray microanalyzer (EDX). The magnetic properties were determined by two methods: electron magnetic resonance (EMR) and vibrating sample magetometer (VSM). We found that cobalt content in the annealed samples reveals an additional magnetic phase transition at lower temperatures.

**Keywords:** Ni-Co-Mn-In ribbons; Heusler alloys; EMR; transformation; annealing

#### **1. Introduction**

The Ni-Mn-based alloys are a very attractive group of multifunctional materials due to their unusual properties and potential application in different areas of life, e.g., in medicine, spintronics, or environmentally friendly refrigeration. Moreover, the physical properties of these materials, such as their magnetocaloric properties or magnetic-shape memory, strongly depend on the atomic ordering and the location of atoms in the crystal structure. Even slight changes in lattice site occupation result in noticeable changes in both magnetic and structural properties of these materials. Such modification can be done in two ways, firstly by substitution of appropriate atoms, as in the case of Ni-Co-Mn-In, which was obtained by substitution of Ni atoms in the ternary alloy with Co [1–4]. The second, simpler method is by applying well-defined heat treatment to the base alloy [5,6]. In the literature, one can find various examples of heat-treating, using Ni-Mn-based alloys, such as annealing with slow cooling or quenching [1,5,7,8]. The heat-treating duration and temperature also vary, but generally, in the ribbons, this time is relatively short compared to bulk samples [9].

The fabrication of ribbons by the melt-spinning method leads to stresses as a result of the rapid solidification process. The annealing process leads to relaxation [10,11], i.e., the removal of such stresses, which are associated with a change in the distance between manganese atoms. This, in turn, may cause changes in the magnetic interactions between these atoms [12], and thus changes in their magnetic properties.

The annealing parameters were chosen on the basis of experiments described in literature [1,5,7,11,13]. The order–disorder transition (ODT), which for the Ni-Mn-In alloy type takes place at 950 K [13], plays a very important role in the selection of the annealing

**Citation:** Dubiel, Ł.; Wal, A.; Stefaniuk, I.; Zywczak, A.; Potera, P.; ˙ Maziarz, W. Cobalt Content Effect on the Magnetic Properties of Ni50−xCoxMn35.5In14.5 Annealed Ribbons. *Materials* **2021**, *14*, 5497. https://doi.org/10.3390/ma14195497

Academic Editor: Hailiang Yu

Received: 26 August 2021 Accepted: 17 September 2021 Published: 23 September 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

temperature. Typically, annealing is carried out at a temperature greater than 950 K. For example, Recarte et al. [7] report that annealing at such temperatures followed by a slow cooling favors the production of a larger fraction of the *L*21 structure than rapid cooling. Although, there are also reports [1] that long-term annealing below the ODT temperature also promotes the production of the *L*21 structure.

Furthermore, there are other annealing methods used in off-stoichiometric Ni-Mn-In and Ni-Mn-Sn alloys [14,15], including annealing in a magnetic field. This procedure leads to the decomposition of the off-stoichiometric alloy into the cubic *L*21 phase and a second phase with lower symmetry [14–17], i.e., ferromagnetic Ni50Mn25X25, (X = In, Sn), precipitates are embedded in a Ni50Mn50 non-magnetic matrix.

In this paper, we investigated annealed off-stoichiometric Ni-Co-Mn-In ribbons with three different compositions. Based on electron magnetic resonance (EMR) spectra and magnetization data, we discussed in detail the magnetic properties of these alloys and their dependence on the cobalt content in the alloy.

#### **2. Materials and Methods**

Three alloys Ni50−xCoxMn35.5In14.5 (x = 0, 3, 5) were prepared by the melt-spinning method according to the procedure, which was described in references [2]. Before measurements the samples were annealed at 1173 K for 30 min and then slowly cooled in a furnace for 12 h. Samples with different cobalt content were labeled NC0MI, NC3MI, and NC5MI for 0, 3, and 5 at.% of cobalt, respectively.

Morphology and composition of samples were inexamined using Tescan Vega 3 scanning electron microscope (TESCAN ORSAY HOLDING, Brno, Czech Republic) equipped with the EDX Bruker Quantax microanalysis system.

The electron magnetic resonance measurements were performed using a Bruker ELEXYS E580 spectrometer (Karlsruhe, Germany) equipped with the Bruker liquid N gas flow cryostat with the 41131 VT digital controller (Bruker Analytische Messtechnik, Rheinstetten, Germany), within the temperature range 110 K ≤ *T* ≤ 450 K. The EMR signal was registered with the following parameters, a 100 kHz modulation and amplitude of 0.1 mT. The microwave power was 23.77 mW for NC0MI, NC5MI and 15.00 mW for NC3MI, respectively. All EMR measurements were performed in the X-band (9.44 GHz). The resonance position was calibrated by DPPH.

The magnetization results were obtained by the use of LakeShore 7407 vibrating sample magnetometer (VSM, Westerville, OH, USA). The field cooled (FC) and the field heated (FH) curves were registered at *μ*0*H* = 0.01 T and *μ*0*H* = 0.1 T in the temperature range of 85 K ≤ *T* ≤ 450 K.

#### **3. Results and Discussion**

#### *3.1. Morphology and Composition*

Figure 1 shows the SEM images of surfaces for melt-spun and annealed (1173 K) NC3MI and NC5MI at room temperature (RT). The all melt-spun ribbons (Figure 1a,c) exhibit a granular surface topology. The grain size is not uniform and grains are randomly distributed throughout the area of ribbons. Although the ribbon is obtained from polycrystalline ingots, the columnar crystal directions in the ribbon are preferentially oriented perpendicular to the ribbon plane [18,19].

The surface morphology of the ribbons after annealing has changed significantly. Additionally there is also a difference in the morphology of the annealed ribbons with different cobalt content.

As shown in Figure 1b, the grains in NC3MI become more uniform after annealing and the largest grains have disappeared. In addition to grains, on the surface of the ribbons one can see the second phase in stripes form. EDS analysis showed that the second phase composition differs from the matrix composition by a significant increase in cobalt concentration (11.35 at.%), at the expense of manganese (29.65 at.%) and indium (11.76 at.%). Figure 1d presents the SEM image of the surface of NC5MI. The morphology of this sample

is very rich and significantly different from the two previous alloys. First of all, the increase in grain size after annealing is clearly visible. In addition to the second phase in the striped form, one can see the eutectoid form on the surface of this ribbon. Similarly to the NC3MI ribbon, the second phase composition is different from the matrix composition.

**Figure 1.** Microstructure of Ni50−xCoxMn35.5In14.5 ribbons at RT: Ni47Co3Mn35.5In14.5 as-spun (**a**) and annealed (**b**) ribbons; Ni45Co5Mn35.5In14.5 as-spun (**c**) and annealed (**d**) ribbons.

#### *3.2. Magnetic Properties*

The magnetic properties of the Ni50−xCoxMn35.5In14.5 ribbons were studied as a function of temperature. The magnetization curves *M*(*T*) obtained for the samples with different content of Co are presented in Figure 2. The FH and FC magnetization curves for NC3MI and NC5MI were obtained in a field 0.01 T. For NC0MI, magnetization cures were registered at 0.1 T due to a very weak signal.

The nature of the *M*(*T*) curves for annealed ribbons differs significantly from nonannealed ribbons [20–22]. The *M*(*T*) of the NC5MI sample displayed in Figure 2f is similar to magnetization of the ribbon before annealing [21]. This magnetization curve does not exhibit any thermal hysteresis and has a clear shift toward higher temperatures relative to non-annealed alloy [20,21].

Also, for the two remaining annealed samples, NC0MI and NC3MI, no clear temperature hysteresis is visible, and the behavior of *M*(*T*), especially for NC0MI, is unusual. In order to better present phase transitions and to determine precisely the temperatures of these transitions, one should consider the temperature dependence of the derivative of magnetization *dM*/*dT* (Figure 2a–c). At high temperature ranges, each sample exhibits a magnetic phase transition from paramagnetism to ferromagnetism with Curie temperatures equal to 300 K, 363 K, and 392 K for NC0MI, NC3MI, and NC5MI, respectively. However, the magnetization below *TC* suggests that these materials are multi-phased. Especially for NC0MI, there is a visible narrow region when magnetization is saturated, but with the further reduction of temperature, the magnetization starts to increase again around 225 K. Similar magnetization curves were observed for Ni50.3Mn35.3In14.4 alloy produced by annealing under a magnetic field [15] and for Ni2.14Mn0.55Sb1.3 after homogenization [23]. Considering the derivative of *M* versus *T* (Figure 2a), one can see that at around 150 K the next magnetic ordering temperature occurs. The character of this magnetic phase transition strongly depends on Co content. With increasing of cobalt concentration, the

intensity of the low temperature peak in the *dM*/*dT* value gradually decreases. For a sample containing 3 at.% of cobalt this peak still exits, but its minimum is much less visible than for NC0MI. In the Figure 2c corresponding to the NC5MI sample the minimum in the low temperature region almost disappears.

**Figure 2.** FH and FC magnetization curves as a function of temperature of Ni50−xCoxMn35.5In14.5 annealed ribbons: Ni50Mn35.5In14.5 (**d**), Ni47Co3Mn35.5In14.5 (**e**) and Ni45Co5Mn35.5In14.5 (**f**). The upper panel contains the corresponding *dM*/*dT* curves (**a**–**c**).

Such behavior of the *M*(*T*) curve may suggest that there are regions of different magnetic properties in the sample, which are revealed during the temperature change. Most likely, these are antiferromagnetic and ferromagnetic regions [14], the mutual volume ratio of which changes during temperature changes. Figure 2a shows that around 150 K, the material undergoes a phase change followed by a further increase in magnetization, meaning that the ferromagnetic region grows at the expense of another one.

Comparing this nature of the change in magnetization as a function of temperature with the results of measurement of the magnetization of a sample with the same composition but not subjected to annealing, it can be concluded that the annealing causes a change in the structure of the Heusler alloy. The authors of [14] report on the phase separation phenomenon in the Ni50Mn45In5 alloy during annealing in the presence of a magnetic field. In our case, annealing took place without an external field and the composition of the sample was different, containing more indium and less manganese. It seems, however, that the annealing process could also in our case cause a change in the structure, i.e., precipitation of various phases of the alloy. Indirect evidence for this may be the composition measurements carried out for the surfaces of the annealed samples, which reveal a different percentage of elements in relation to the unheated samples [20]. The comparison of Figure 2a–c shows that increasing the cobalt concentration causes homogeneity of these areas, i.e., the structural heterogeneity in the volume of the alloy decreases.

Another explanation of the observed properties can be based on [23], the authors of which report a similar relationship *M*(*T*) of Ni2.14Mn0.55Sb1.31 samples subjected to annealing at the temperature of 850 ◦C. They assume that the crystal structure of the samples is the same in the whole volume, but the elements that compose it may be in different places in different areas. This can lead to magnetic interactions of the FM and AFM type, which compete with each other at different temperatures, resulting in the observed magnetization.

Figure 3 displays some selected EMR spectra for Ni50−xCoxMn35.5In14.5 ribbons with different concentrations of Co atoms. The spectra were registered within the temperature range of 110 K ≤ *T* ≤ 450 K. In general, the annealing process does not drastically change the character of the EMR signal registered for Ni50−xCoxMn35.5In14.5 [20,21]. At the highest temperatures, the main line has an unchanged asymmetric character. Such an observation confirms that the ribbons before annealing were already partially ordered materials.

**Figure 3.** Selected EMR spectra of derivative of the resonance absorption (*dP*/*dB*) for Ni50−xCox Mn35.5In14.5 ribbons: Ni50Mn35.5In14.5r (**a**), Ni47Co3Mn35.5In14.5 (**b**) and Ni45Co5Mn35.5In14.5 (**c**) registered during a cooling process. Spectra are vertically shifted for clarity. The example of the LFMA signal is marked in red.

At the highest temperatures the EMR spectra of NC0MI ribbon include two separated Dyson lines: very weak line 1 (*BR* ∼ 330 mT) and line 2 (*BR* ∼ 200 mT). Within the temperature range 400 K ≤ *T* ≤ 450 K, the amplitude of the line 1 is so weak that its contribution to the EMR signal can be neglected. Below 400 K, the amplitude of line 1 gradually increases, and at a temperature of about 345 K, line 1 reaches a height equal line 2. With the further reduction in temperature, the intensity of line 1 increases and this line shifts towards weaker fields, and in 300 K it overlaps with line 2. For this sample, the mainline (line 1) is similar in nature to the paramagnetic lines observed for the Ni-Mn-In ribbons in previous work [20,24].

The spectra of the sample NC3MI are presented in Figure 3b. In the temperature range 426 K ≤ *T* ≤ 445 K, a single strong, asymmetric line is observed. This is a typical paramagnetic line. Below 426 K EMR line starts broadening and shifts toward a low-field region. Figure 3c shows the EMR spectra for NC5MI registered within the temperature range of

270 K ≤ *T* ≤ 450 K. At the highest temperatures in the EMR spectrum, a single asymmetrical Dysonian line (line 1) is visible. In the range of temperatures 430 K ≤ *T* ≤ 450 K, the position of the line slightly changes. During cooling of the sample, below 425 K the main line shifts towards the low field region, and, additionally, in the EMR spectrum around 340 mT a much weaker symmetrical line appears (line 2). As the temperature decreases these two lines become better separated.

The resonance conditions below the magnetic phase transition temperature in ferromagnetic materials can be considered within the classical theory of ferromagnetic resonance proposed by Kittel [25]. The thickness of the ribbon (parallel to *x*-axis) is far smaller than other sample dimensions and we can treat it, in the first approximation, as a plane surface, where the magnetic field is parallel to it. This model corresponds to a special case of the Kittel equation, where the demagnetization factors are equal to *Nx* = 1 and *Ny* = *Nz* = 0, and the Kittel equation will be simplified considerably (SI units):

$$
\omega\_0 = \gamma \left[ (\mu\_0 M + B) B \right]^{1/2} \prime \tag{1}
$$

where *ω*0 is the microwave angular frequency (*ω*0 = 2*π f*0), *γ* is the gyromagnetic ratio, *B* means the static magnetic field, *μ*0 stands for the permeability, and *M* denotes the saturation magnetization. A Typical continuous wave EMR experiment has a fixed microwave frequency, solving the Equation (1) for *M*, assuming *B* as the value of the magnetic field corresponding to the resonance condition.

Table 1 lists the *M* values calculated based on Equation (1) and determined from the hysteresis loops for particular temperatures. The latter value of magnetization *M* for each temperature was obtained from *M*(*H*) dependence presented in Figure 4. As can be seen for the NC3MI ribbon, the magnetization value *M* determined from Equation (1) is slightly higher than that from the figure. This difference can be related to the fact that EMR probes only the surface of a sample (∼1 μm), while the VSM measurements provide information about bulk magnetization. Thus, in the EMR results, the outer material layer is taken into account, where the interaction of a ferromagnetic nature is dominant, while VSM measures the magnetization of the entire sample volume, thus the total magnetization may be reduced as a result of the presence of a different magnetic ordering.

**Label** *T* **(K)** *M* **(kA/m)** *M∗* **(kA/m)** NC0MI 300 147 26 100 608 157 NC3MI 350 330 242 300 553 416 100 879 694

**Table 1.** The magnetization value determined from the solution of the Kittel equation (*M*) and from hysteresis loops (*M*∗) for NC0MI and NC3MI samples.

At temperatures below *TC*, in the EMR spectra registered for each sample, an additional signal appears, known in the literature as low-field microwave absorption (LFMA). Generally, the LFMA is observed for soft ferromagnetic materials, such as manganites, amorphous ribbons, magnetic thin films, magnetic nanoparticles [26–30]. In addition, LFMA is also used to detect ferromagnetism and to determine *TC*. Furthermore, the absence of LFMA in the soft ferromagnetic sample is a good indicaton of existence of superparamagnetic state [29,30].

**Figure 4.** The field dependence of magnetization for Ni50Mn35.5In14.5 (**a**) and Ni47Co3Mn35.5In14.5 (**b**) ribbons at the selected temperatures.

In our prior works, we observed LFMA in the Ni-Mn-In and Ni-Co-Mn-In ribbons [22,24] but the behaviour of LFMA in annealed ribbons is more complex. Below *TC* we observe only the fragment of the LFMA signal, but this signal has the same phase as the main line. With the further reduction in temperature, the LFMA signal changes their phase. The change in the signal phase for the sample without cobalt occurs at a temperature of approximately 291 K, for the sample with the cobalt content of 3 at.%, this temperature is higher, approx. 351 K. For the sample with the highest cobalt content (5 at.%), the phase change temperature is approximately 370 K. An especially interesting behavior of the LFMA can be observed for the NC3MI ribbon. The Curie temperature determined from VSM data for sample NC3MI is equal to 364 K (see Figure 2b), while the value of *TC* determined from the temperature dependence of the integrated intensity of EMR line is equal 421 K (see Table 2). The difference in the value of Curie temperature obtained from VSM and EMR is a consequence of the different natures of both techniques. The EMR is sensitive to changes in the environment of the magnetic ions, so it is a good method for detecting local changes in properties of materials [21]. For conducting samples, due to an existing skin depth phenomenon, and the local character of the EMR method, one can notice a discrepancy between results determined from EMR data and VSM technique. In our prior measurements, however, for as-spun Ni45Co5Mn35.5In14.5 [21] and annealed Ni50Mn35.5In14.5 [24], only a slight difference in *TC* values determined by using the EMR and VSM technique was observed, although for those samples the skin depth effect also exists. The difference in the Curie temperature values determined in both methods for

samples described in this report may be explained by the annealing process, which caused the inhomogeneity of the samples. Ferromagnetic structures may have formed on the sample surface, which is manifested by the Curie 421 K temperature determined by the EMR method and the differences in the magnetization values calculated using the EMR and VSM techniques and those listed in Table 1. The sample taken as a whole, neglecting the thin layer near the surface, remains paramagnetic at this temperature, as indicated by *M*(*T*) and *M*(*H*) dependences. Ferromagnetism in the macro scale is revealed at a lower temperature, about 360 K, which is visible both in the magnetization data and in the anomalous behaviour of the EMR line, suggesting the occurrence of a phase transition.

**Table 2.** The corresponding EMR results of NC0MI, NC3MI, and NC5MI samples and comparison with literature data. *TC* are determined from the temperature dependence of integral intensity, *a* and *b* parameters are determined from temperature dependence of the EMR linewidth.

\* annealed at 600 K; \*\* non-annealed.

Figure 5 exhibits the temperature dependence of linewidth Δ*B*(*T*) for NC0MI and NC5MI. In the temperature range of 280 K ≤ *T* ≤ 450 K the EMR spectra for NC0MI (Figure 5a) consists of two lines. In the whole temperature range, the Δ*B* of the line 2 varies slightly between 30–40 mT. In contrast, Δ*B* of line 1 in the temperature range of 330 K ≤ *T* ≤ 400 K decreases linearly, and then, after approaching Curie temperature, the linewidth starts to increase. This behaviour of Δ*B*(*T*) around *TC* is typical for ferromagnetic metals [31]. The high temperature EMR signal for NC5MI within temperature range 395 K ≤ *T* ≤ 450 K can be also divided into two lines, but the linewidth for NC5MI (Figure 5b) has a different character than the temperature dependence for NC0MI. The main line, which is labelled 1, shows linear behaviour of Δ*B*(*T*) only in a narrow range of temperature (430 K ≤ *T* ≤ 450 K) and its value is about 30 mT. Below these temperatures, the Δ*B* of both line increases. For NC3MI, no such relationship was determined due to the unregularly nature of changes in linewidth as a function of temperature.

The linear region of Δ*B*(*T*) observed in the temperature dependence for both samples, can be expressed by the Korringa relaxation model [32]:

$$
\Delta B(T) = a + b \cdot T,\tag{2}
$$

where *a* = Δ*B*0 is the residual width and *b* is the Korringa rate.

The value of the last parameter *b*, describing the slope of Δ*B*(*T*), depends mainly on concentration of 3*d* elements, i.e., the *b* value increases when the number of 3*d* electrons decreases [33]. For NC0MI, the value of *b* is about 10 G/K (see Tabel 2) and this value is typically observed for transition metal EMR [34]. A similar value of this parameter was determined for the in-situ ribbon [20]. For the linear part of line 1 of NC5MI, the value of b is smaller and equal to 1 G/K. A similar value was observed in YBaMn2O6 [35], where the bottleneck scenario for the spin relaxation of the manganese ion via mobile *eg* electron was noticed.

The main values determined from the EMR results, such as *TC*, *a*, and *b* are collected in Table 2. Moreover, the table contains the EMR parameter values for in-situ [21,22] and another annealed ribbon [24] presented in our earlier works. The smallest changes in the EMR parameter values were observed for the sample without cobalt. However, in the paramagnetic regime, a weak additional line can be observed, which can be considered a singularity for Ni-Mn-In ribbons. Other parameters, including the Curie temperature, differ only by 2 K, which shows that the annealing parameters are not native, and the annealing itself significantly shifts the Curie temperature, i.e., for the in-situ ribbon, this temperature is 267 K [20] and it overlaps with the martensitic transformation temperature. It looks different in the case of NC3MI and NC5MI ribbons. For the sample with 3 at.% of cobalt the temperature dependence of Δ(*B*) has an irregular character and cannot be described using the Korringa relaxation model, while for as-cast ribbon with the same chemical composition, one can observe a typical Korringa relaxation process [22]. Furthermore, the Curie temperature, determined from the temperature dependence of the integral intensity, increases significantly relative to the non-annealed sample (*TC* = 340 K) [24] and differs from the Curie temperature for the NC3MI sample, determined from the temperature dependence of magnetization. In the case of the NC5MI ribbon, unlike the NC0MI ribbon and the non-annealed NC3MI ribbon, the width of the paramagnetic line (line 1) does not change significantly with temperature changes (above 430 K) and oscillates around the value ∼30 mT, and the parameters of the linear fit of the fragment Δ*B*(*T*) above *TC*, equals *a* ≈ 0 mT, and *b* = 1 G/K. Such parameter values and the shape of Δ*B*(*T*) indicate that the temperature dependence of the line width for the annealed NC5MI ribbon cannot be described by the Korring relaxation model.

**Figure 5.** Temperature dependence of the EMR linewidth Δ*B* for Ni50Mn35.5In14.5r (**a**) and Ni45Co5Mn35.5In14.5 (**b**) ribbons. The solid lines represent the best fit of the linewidth above *TC*, using the expression (1).

#### **4. Conclusions**

The effect of cobalt content of annealed Ni50−xCoxMn35.5In14.5 for (x = 0, 3, 5) is investigated through scanning electron microscopy, electron magnetic resonance, and vibrating sample magnetometer. The main conclusions can be listed as follows:

**Author Contributions:** Conceptualization, Ł.D.; methodology, Ł.D. and A.W.; validation, I.S. and A.Z.; formal analysis, Ł.D. and A.W.; investigation, Ł.D., I.S., A. ˙ Z., P.P. and W.M.; writing—original ˙ draft preparation, Ł.D.; writing—review and editing, Ł.D. and A.W. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.