Experimental study of the liquidus surface in the V-rich portion of the V-Si-B system

Journal Pre-proof Experimental study of the liquidus surface in the V-rich portion of the V–Si–B system G. Hasemann PII:

S0925-8388(20)30206-1

DOI:

https://doi.org/10.1016/j.jallcom.2020.153843

Reference:

JALCOM 153843

To appear in:

Journal of Alloys and Compounds

Received Date: 27 October 2019 Revised Date:

10 January 2020

Accepted Date: 13 January 2020

Please cite this article as: G. Hasemann, Experimental study of the liquidus surface in the V-rich portion of the V–Si–B system, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/ j.jallcom.2020.153843. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

Credit Author Statement

G. Hasemann: Writing - original draft & review, Investigation

Experimental Study of the Liquidus Surface in the V-rich portion of the V-Si-B System G. Hasemann1,2

1

Otto-von-Guericke University Magdeburg, Institute of Materials and Joining Technology, Universitätsplatz 2, 39106 Magdeburg, Germany

2

Forschungszentrum Jülich GmbH, Institute of Energy and Climate Research, Microstructure and Properties of Materials (IEK-2), Leo-Brand-Str. 1, 52425 Jülich, Germany Abstract

The solidification behavior of arc-melted alloys in the V-rich portion of the ternary VSi-B alloy system has been experimentally investigated. A detailed microstructure analysis of several as-cast alloys based on SEM observations, EDS/WDS and EBSD measurements was carried out. As a result, different primary solidification reactions in the V-rich V-Si-B system were identified. On the two binary boundary systems V-Si and V-B those reactions were attributed to VSS, V3Si, V5Si3, V3B2 and VB. Furthermore, a primary solidification reaction within the ternary system was attributed to V5SiB2. As a result, the new data allows a reconstruction of the liquidus surface of the V-Si-B phase diagram.

Keywords: Intermetallics, Transition metal alloys and compounds, Microstructures, Scanning electron microscopy

1. Introduction

The ternary V-Si-B alloy system has recently gained new attentions by the work of the Krüger group [1–4]. Due to the similarities between the V-Si-B system and its wellinvestigated Mo-Si-B neighboring system strong similarities in the alloy design strategies were determined, especially the alloy design concepts for powder metallurgical (PM) processed Mo-Si-B alloys were transferred and applied on V-Si-B materials [2]. 1

At present, however, the V-Si-B system is purely investigated as compared to its MoSi-B neighbor system. Early investigations on the V-Si-B phase diagram can be referred to Kudielka and Nowotny [5]. Their investigations revealed that the V5SiB2 (T2) phase has the same Si/B ration as compared to the Mo-Si-B system and postulated a first isothermal section at 1450 °C. Depending on the alloys’ composition, they identified phase equilibria between V5SiB2-VB, V5SiB2-D88 and V5SiB2-V5Si3-D88, where the D88 phase is a V5Si(3-x)Bx phase and represents a B-stabilized V5Si3 (T1) phase [5,6]. A second approach on the V-rich V-Si-B system had been reported by Nunes et al. [6], who prepared various alloys via arc-melting and annealed them to investigate the phase equilibria. Their results led to an isothermal section at 1600 °C. Besides isothermal sections of ternary phase diagrams for studying and understanding phase equilibria at certain compositions and/or temperatures, the liquidus projection of a certain system is often a very useful tool to investigate the cooling behavior and solidification paths of alloys. To know about the solidification path and reaction sequences may further lead to understand phase transformations during casting processes. Thus, understanding solidification processes of complex multiphase alloys helps to further optimize the materials processing. While there is an abundance of literature information about the liquidus surface and microstructure evolution in the MoSi-B system [7–16], only little is known about the solidification behavior of the V-based system. First investigations on the as-cast microstructures of ternary V-Si-B alloys were obtained by Reis et al. [17]. Their work provides information of possible primary phase formations during cooling from the liquid phase of only seven V-Si-B alloys. All alloys were located within the V-rich portion of the phase diagram and the solid phases were identified by using wavelength dispersive spectroscopy (WDS). Recently Da Silva et al. [18] thermodynamically modeled the V-Si-B system by using the CALPHAD method. They optimized their calculations by recalculation the binary VSi, V-B and Si-B phase diagrams and used data from recent assessments of all three systems [19–22]. Da Silva et al.’s [18] thermodynamic modeling involved a calculation of an isothermal section at 1600 °C and a first liquidus projection of the ternary V-Si-B system and supported their calculations with a few experimental information. This work allows a first insight into the possible invariant reactions of the complete V-Si-B system including calculated compositions and temperatures of those four-phase reactions. 2

To investigate the liquidus projection, the binary boundary systems should also be checked, respectively. Comprehensive information on the binary phase diagrams of V-Si and V-B are given by Smith [23] and Spear et al. [24]. Those data had been critically reinvestigated recently [21,25] and were further refined and optimized with the help of thermodynamic calculations using the CALPHAD method [18,19]. Thus, the solidification paths of both V-rich binary systems are summarized in Table 1, involving eutectic and peritectic reactions, reaction temperatures and liquid compositions.

Tab. 1 Invariant reactions in the binary equilibrium diagrams of V-Si and V-B Type

Invariant Reaction

Liquid Composition

Temperature

Reference

e

L ↔ VSS + V3Si

12.6 at.% Si

1837 °C

[25,26]

e

L ↔ V3Si + V5Si3

27.6 at.% Si

1895 °C

[25]

e

L ↔ VSS + V3B2

11.3 at.% B

1732 °C

[18]

p

L + VB ↔ V3B2

18.7 at.% B

1905 °C

[18]

e: eutectic reaction; p: peritectic reaction

Besides experimental investigations by De Lima-Kühn et al. [25], the binary VSS-V3Si eutectic had been well-studied by Bei et al. [26], who reported a fully eutectic microstructure at 12.6 at.% Si. Values slightly smaller than suggested by the binary phase diagram [23] were further corroborated by Bewlay and Chang [27,28] and Henschall et al. [29]. Hence, at least the solidification behavior of the binary V-Si system is well-investigated. The recent experiments by De Lima-Kühn et al. [25] further suggest, that the second V3Si-V5Si3 eutectic reaction tends to be also slightly leaner in its Si concentration, namely 27.6 at.%, as compared to the V-Si phase diagram proposed by Smith [23]. The binary V-B system is, however, less well-investigated. Again, experimental investigation by De Lima et al. [21] reported a slightly lower B-concentration for the binary VSS-V3B2 eutectic reaction at 12 at.%, as compared to 13 at.% B of the V-B phase diagram [24]. Recent information on the V-B system are based on optimized thermodynamic calculations [18]. The present work studies the liquid - solid phase transformation in the V-rich portion of the V-Si-B system experimentally. The solidification behavior of various V-Si-B alloys 3

is investigated in terms of their phase formation and microstructural evolution during cooling down from the liquid state. The experiments identified five invariant reactions in the V-rich part of the ternary system and allow drawing a liquidus surface based on the present experimental results.

2. Experimental procedures

To carry out the experiments various V-Si-B alloys were prepared via conventional arc-melting in an Ar atmosphere. Prior to melting, the alloy concentrations were carefully weight in using high purity elemental chips or flakes of V (99.7 %), Si (99.99 %) and B (99.0 %). The alloys were flipped and re-melted at least five times to ensure good homogeneity. The weight loss after melting was ≤ 1 % indicating almost no deviations from the nominal, weight-in composition. The chemical alloy compositions were verified by using inductively coupled plasma optical emission spectroscopy (ICP-OES). Table 2 summarizes the as-cast alloys investigated in the present study providing information on their individual chemical compositions and microstructure analysis. Samples for metallographic preparation were embedded in a hot mounting epoxy (Struers Poly Fast) and subsequently grinded followed by a finish with 3 µm and 1 µm diamond suspension and colloidal silica. The microstructural observations were carried out by using a Zeiss Merlin and Zeiss Supra 50 VP scanning electron microscope (SEM). The SEM images were typically obtained in the backscattered electron (BSE) mode. For the phase identification wavelength-dispersive X-ray spectroscopy (WDS, Oxford Instruments) and electron backscatter diffraction (EBSD, Oxford Instruments) was used. For WDS, pure V, Si and B were used as standards and a window-free energy dispersive X-ray spectroscopy system (EDS, Oxford Instruments) was used to double-check the results. Quantitative boron detection in the V-Si-B system is less problematic as compared to the Mo-Si-B system, since V (Lα-line at 0.511 eV) and B lines do not overlay as it is known for the Mo Mζ-line (0.193 eV) and the B Kα-line (0.183 eV).

4

Tab. 2 Chemical analysis of as-cast V-Si-B alloys obtained by ICP-OES and their microstructure characteristics identified by WDS and EBSD No.

Nominal composition

Si [at.%]

B [at.%]

Primary Phases

1

As-cast Phases

V-7Si

7.4

-

VSS

VSS, V3Si

2

V-5Si-5B

5.2

4.7

VSS

VSS, V3B2, V5SiB2

3

V-8.5Si-6B

8.7

5.9

VSS

VSS, V5SiB2

4

V-9Si-5B

9.3

5.0

VSS

VSS, V3Si, V5SiB2

5

V-6.6Si-7.3B

6.8

7.0

VSS

VSS, V3B2, V5SiB2

6

V-10Si-2.5B

9.7

2.4

VSS

VSS, V3Si, V5SiB2

7

V-11Si-1B

10.8

0.9

VSS

VSS, V3Si, V5SiB2

8

V-3Si-15B

3.1

14.7

V3B2

V3B2, VSS

9

V-8Si-7.5B

8.2

7.4

V3B2

V3B2, VSS, V5SiB2

10

V-2Si-12B

2.1

11.5

V3B2

V3B2, VSS

11

V-4Si-12B

4.1

11.5

V3B2

V3B2, VSS, V5SiB2

12

V-1Si-17B

1.0

17.3

V3B2

V3B2, VSS

13

V-8.8Si-12.4B

8.8

12.4

V3B2

V3B2, VSS, V3Si, V5SiB2

14

V-5Si-9B

5.2

8.3

V3B2

V3B2, VSS, V5SiB2

15

V-8Si-10B

8.1

9.8

V3B2

V3B2, VSS, V5SiB2

16

V-1Si-10B*

1.1

9.9

V3B2

V3B2, VSS (bin. eutectic)

17

V-11B

-

11.5

V3B2

V3B2, VSS

18

V-11Si-3.5B

11.1

3.5

V3Si

V3Si, VSS, V5SiB2

19

V-12Si-7B

12.1

6.5

V3Si

V3Si, VSS, V5SiB2

20

V-15Si-6B

14.4

6.2

V3Si

V3Si, VSS, V5SiB2 5

21

V-16.5Si-3.5B

16.7

3.5

V3Si

V3Si, VSS, V5SiB2

22

V-17Si-6B

17.6

5.9

V3Si

V3Si, VSS, V5SiB2

23

V-20Si-6.5B

20.4

6.3

V3Si

V3Si, VSS, V5SiB2

24

V-9Si-7B

8.4

7.1

V5SiB2

V5SiB2, V3Si, VSS

25

V-10Si-10B

10.3

8.7

V5SiB2

V5SiB2, V3Si, VSS

26

V-10Si-7.5B

10.3

7.3

V5SiB2

V5SiB2, V3Si, VSS

27

V-17Si-10B

17.1

9.6

V5SiB2

V5SiB2, V3Si, VSS

28

V-13Si-11B

13.1

11.1

V5SiB2

V5SiB2, V3Si, VSS

29

V-23Si-6B

23.4

5.5

V5SiB2

V5SiB2, V3Si, V5Si3

30

V-3Si-17B

2.9

17.6

VB

VB, V3B2, VSS

31

V-6.5Si-13B

6.7

12.8

VB

VB, V3B2, VSS, V5SiB2

32

V-9Si-13B

9.0

13.0

VB

VB, VSS, V5SiB2, V3Si

33

V-17Si-11B

17.0

10.9

VB

VB, VSS, V5SiB2, V3Si

34

V-20Si-10B*

20.2

10.5

VB

VB, V5SiB2, V3Si

35

V-25Si-8B*

25.6

8.9

VB

VB, V5Si3, V5SiB2, V3Si

36

V-27Si-8B*

25.8

8.1

VB

VB, V5Si3, V5SiB2, V3Si

37

V-27Si-9B*

26.9

10.2

VB

VB, V5Si3, V5SiB2, V3Si

38

V-31Si-7B

31.8

7.7

VB

VB, V5Si3, V5SiB2, V3Si

39

V-27Si-1B*

27.1

0.9

V5Si3

V5Si3, V3Si, V5SiB2

40

V-26Si-3B*

25.5

3.4

V5Si3

V5Si3, V3Si, V5SiB2

41

V-25Si-6B

24.2

6.0

V5Si3

V5Si3, V3Si, V5SiB2

42

V-24Si-5B

24.1

4.6

V5Si3

V5Si3, V3Si, V5SiB2

* Alloy split and crack while cooling down or heating up with the arc

6

3. Results

In the following sections the specific primary solidification areas identified within the V-rich portion of the V-Si-B system will be investigated in detail. To carry out this work, previous experiments by De Lima [30] had been used as orientations to identify a certain area of interest within the large V-rich portion of the ternary V-Si-B diagram. In the present study the experimental work had been extended and different alloy compositions were chosen to investigate their as-cast microstructure and to define different primary solidification areas. The experimental alloy compositions are summarized in Table 2 and their microstructural evolution will be discussed in more detail.

3.1 Alloys with VSS primary phases Microstructures of alloys solidifying within the primary VSS field are exemplarily shown in Figure 1. Alloy V-5Si-5B (#2) solidifies via two binary eutectic valleys of VSS-V3B2 and VSS-V5SiB2. Due to electron channeling effect of the V3B2 and the V5SiB2 those twophase eutectics are almost impossible to distinguish by using SEM-BSE images. However, the existence of the small fractions of the VSS-V3B2 was recently shown by additional EBSD measurements [31]. In contrast, in alloy V-8.5Si-6B (#3) only the VSSV5SiB2 two-phase eutectic is observed after primary VSS dendrites had been formed during solidification. Alloy V-9Si-5B (#4) shows a relatively high volume fraction of the primary VSS phase. However, the alloy is believed to be close to a class I (E-type) ternary eutectic reaction L ↔ VSS + V5SiB2 + V3Si (E′). Since only few two-phase eutectic VSS-V5SiB2 regions can be found, the cooling path of alloys V-9Si-5B (#4) seems to be strongly influenced by kinetic and undercooling effects. Strong undercooling can lead to eutectic solidification of hypo- or hypereutectic two-phase eutectics and might be a plausible explanation of some undercooling effects observed for various V-Si-B alloys of this present study [32,33]. The finding of a ternary eutectic reaction within the three-phase field of VSS, V5SiB2 and V3Si is, however, very promising for possible future investigations and application of V-rich V-Si-B alloys as low density high temperature alloys. 7

Fig. 1 VSS primary solidifying alloys obtained using SEM-BSE images: a) V-5Si-5B (#2); b) V-8.5Si-6B (#3) and c) V-9Si-5B (#4). The ternary eutectic observed in alloy V9Si-5B (#4) is shown in more detail in d).

3.2 Alloys with V3B2 primary phases Exemplary, the as-cast microstructures of V3B2 primary solidifying alloys are shown in Figure 2. In alloy V-4Si-12B (#11) and V-1Si-17B (#12) a two-phase VSS-V3B2 type eutectic is observed after V3B2 had formed primarily. Thus, the present observations give evidence for a mono-variant line between the binary eutectic VSS-V3B2 point and alloy V-5Si-9B (#14). The solidification path of this alloy ends on the VSS-V5SiB2 eutectic line, which could be determined by EBSD measurements. Due to similar electron channeling contrasts between the V3B2 and V5SiB2 phase, the second two-phase eutectic cannot be recognized in the SEM-BSE images, which was also checked via EBSD. The alloy V-8Si-10B (#15) is located on the Si-rich part of the V3B2 primary solidification field. After V3B2 primary solidification, the phase seems to transform in a peritectic-type reaction forming a VSS-V5SiB2 quasi-binary eutectic and negligible small 8

portions of the ternary eutectic. Thus, the line between the two primary solidification fields of V3B2 and V5SiB2 has a peritectic character since a eutectic reaction between these two phases could not be observed and seems not to exist. Due to the present experimental observation, a class II (U-type) reaction of L + V3B2 ↔ VSS + V5SiB2 (II′) is assumed.

Fig. 2 V3B2 primary solidifying alloys obtained using SEM-BSE images: a) V-4Si-12B (#11), b) V-1Si-17B (#12), c) V-5Si-9B (#14) and d) V-8Si-10B (#15)

3.3 Alloys with V3Si primary phases Figure 3 shows three alloys which are attributed to the V3Si primary solidification field in the present work. Alloy V-11Si-3.5B (#18) shows primary solidification of large V3Si phases. However, the solidification path of the alloy ends at the VSS-V5SiB2 eutectic valley, which seems to be quite unusual. Due to rapid solidification via arc-melting the residual melt spontaneously decreases in its Si-concentration and VSS might be formed via secondary 9

phase formation [32]. Since the alloy seems to be very close to the VSS-V3Si eutectic line the rapid Si-depletion in the residual melt overjumps the two-phase eutectic line and after secondary VSS formation the solidification sequence ends at the VSS-V5SiB2 eutectic valley. This explanation can be supported by the specific geometrical shape of the V3Si and VSS phases. Both are round-shaped and in contact with each other. In a closer look, the VSS mostly touches the V3Si phase with concave faces. This would be a strong indicator, that the V3Si phase has formed prior to the VSS phase. Furthermore, the VSS and V3Si exhibit an electron channeling effect (as described earlier for V3B2 and V5SiB2) and are thus, difficult to distinguish – even by using the SEM-BSE contrast mode so that EBSD is recommended. The alloys V-12Si-7B (#19) and V-20Si-6.5B (#23) solidify very close to the V3SiV5SiB2 eutectic valley which is indicated by only minor portions of primary V3Si phase and relatively large areas of two-phase V3Si-V5SiB2 eutectics. Depending on the remaining fraction of liquid, larger (V-12Si-7B (#19)) or smaller sections (V-20Si-6.5B (#23)) of VSS-V5SiB2 eutectic can be observed.

Fig. 3 V3Si primary solidifying alloys obtained using SEM-BSE images: a) V-11Si3.5B (#18), b) V-12Si-7B (#19) and c) V-20Si-6.5B (#23) 10

It has to be mentioned here, that the primary solidification area of the V3Si seems to be strongly affected by undercooling effects and by solidification of secondary phases, which makes it very difficult at the present state, to draw any detailed conclusions on the solidification path of those alloys. Nevertheless, the dimension of the V3Si primary field can be estimated relative to the results of the neighboring primary solidification areas of VSS, V5SiB2 and V5Si3. 3.4 Alloys with V5SiB2 primary phases The primary solidification area of V5SiB2 is determined by the alloys exemplary shown in Figure 4 and their different solidification paths. In alloy V-10Si-10B (#25) large V5SiB2 phases are preceding a two-phase VSS-V5SiB2 eutectic. Additionally, small but coarse sections of a ternary VSS-V5SiB2-V3Si eutectic can be observed. Large V5SiB2 primary phases can also be detected in alloy V-17Si-10B (#27). Upon further cooling a coarse V3Si-V5SiB2 eutectic is forming which is followed by small and extremely fine VSS-V5SiB2 eutectic areas. Due to undercooling effects, the alloys seem to skip the ternary eutectic reaction which might explain the unexpected presence of the fine VSS-V5SiB2 eutectic. Thus, after primary solidification the solidification path seems to be similar with alloy V-12Si-7B (#19) and V-20Si-6.5B (#23). In contrast, alloy V-23Si-6B (#29) is located close to a four-phase reaction which can be

identified

as

a

ternary

eutectic

reaction

of

three

intermetallic

phases,

L ↔ V5SiB2 + V3Si + V5Si3 (E″). According to the small volume fraction of the primary V5SiB2 phase, alloy V-23Si-6B (#29) is believed to be very close to the ternary eutectic point E″.

11

Fig. 4 V5SiB2 primary solidifying alloys obtained using SEM-BSE images: a) V-10Si10B (#25), b) V-17Si-10B (#27) and c) V-23Si-6B (#29). The second ternary eutectic observed in the V-rich V-Si-B system of alloy V-23Si-6B (#29) is shown in more detail in d).

3.5 Alloys with VB primary phases

The primary solidification field of the VB phase represents the largest one within the V-rich portion of the V-Si-B system. It ranges from the peritectic reaction at 18.7 at.% B [18] on the V-B binary system well over 40 at.% Si on the Si-rich side. The microstructure formation of alloy V-3Si-17B (#30) in Figure 5a) on the B-rich side of the ternary V-Si-B system is controlled by a sequence of peritectic transformations. After the VB phase has formed primarily, it transformed into the V3B2 phase. The strong halo-formation in Figure 5a) indicates this transformation as a peritectic one, which is in agreement with the binary V-B system [18,24]. As a consequence, the solidification path of the remaining melt is comparable to alloys solidifying in the primary V3B2 area, i.e. alloys V-1Si-17B (#12) forming a VSS-V3B2 eutectic. 12

The as-cast microstructure of alloy V-9Si-13B (#32) indicates that VB seem also to transform into V5SiB2 in a peritectic manner, Figure 5b). The solidification seems to continue along the VSS-V5SiB2 eutectic valley ending in the ternary eutectic point VSSV5SiB2-V3Si (E′). The V3B2-V5SiB2 mono-variant line seems to be not preferential upon further cooling (as compared to VSS-V5SiB2) which would indicate a weak peritectic character of this reaction line. Figure 5c) shows the microstructure of alloy V-25Si-8B (#35). After primary VB formation the phase seems to transform in a peritectic manner into V5SiB2 and V5Si3, which is in agreement with Da Silva et al. [18] and the U-type reaction L + VB ↔ V5SiB2 + V5Si3. Upon further cooling the solidification path follows the monovariant V5SiB2-V5Si3 line. However, the solidification of alloy V-25Si-8B (#35) does not end with the Si-rich ternary eutectic reaction. Due to undercooling effects and the present microstructure observations of the V-Si-B system it is believed that the remaining melt overshoots the L ↔ V5SiB2 + V3Si + V5Si3 ternary eutectic reaction and, as a consequence, the solidification path ends on the V3Si-V5SiB2 line.

Fig. 5 VB primary solidifying alloys obtained using SEM-BSE images: a) V-3Si-17B (#30), b) V-9Si-13B (#32) and c) V-25Si-8B (#35) 13

3.6 Alloys with V5Si3 primary phases The primary solidification area of the V5Si3 was determined by looking especially on the V3Si-V5Si3 eutectic valley and the second, purely intermetallic ternary eutectic reaction L ↔ V5SiB2 + V3Si + V5Si3 (E″).

Fig. 6 V5Si3 primary solidifying alloys obtained using SEM-BSE images: a) V-27Si-1B (#39) and b) V-24Si-5B (#42). The second ternary eutectic observed in the V-rich V-Si-B system of alloy V-24Si-5B (#42) is shown in more detail in c).

The alloy V-27Si-1B (#39), shown in Figure 6a), is located very close to the V3SiV5Si3 binary eutectic reaction. A coarse binary eutectic microstructure can be observed which is comparable with the binary eutectic alloy V-27.6Si published by De Lima-Kühn et al. [25]. EBSD measurements also revealed the allotropic transformation of the V5Si3 phase form the hexagonal, D88 structured high temperature phase to the tetragonal, D8m structured low temperature phase [23,34]. The retained hexagonal crystal structure

of V5Si3 may emphasize slow transformation kinetics or might be due to fast cooling 14

during the arc-melting process. The latter applies also in the case of undercooling effects which might be the reason for secondary phase formation observed within the as-cast microstructures of many V-Si-B alloys studied in the present work. Furthermore, the alloy V-24Si-5B (#42), which is shown in Figure 6b) is of interest. The alloy is very close to the second ternary eutectic point (E″) of the V-rich liquidus projection. The microstructure consists of only small portions of primary V5Si3 phase. As observed for alloy V-27Si-1B (#39) before, the allotropic configuration of V5Si3 could also be found via EBSD. After primary solidification the cooling paths follows the two-phase V5SiB2-V5Si3 and ends in the purely intermetallic ternary eutectic reaction of L ↔ V3Si + V5SiB2 + V5Si3. The ternary eutectic microstructure is shown in more detail in Figure 6c) and marks a very important invariant point within the liquidus projection.

4. Discussion

A first approach on the V-Si-B liquidus projection had been published recently by Da Silva et al. [18] and is based on CALPHAD calculations. Their results, shown in Figure 7, are evaluated by employing the Alkemade theorem to the postulated liquidus surface. The Alkemade theorem defines the direction of decreasing temperature along a mono-variant reaction line on a liquidus surface [35,36]. By using Alkemade lines, it is possible to identify local maxima (more precisely saddle points) along eutectic valleys. Hence, the theorem can be used to predict solidification paths of various alloys in a ternary system and is useful to determine possible invariant ternary reactions. It should be noted that the theorem in its original form is specifically defined for line compounds with no (or negligible) solid solubility. However, many systems show certain solubility ranges of intermetallic phases. Not only binary phases but also ternary phase (for example the Mo5SiB2 phase in the Mo-Si-B system [13,37]) can exhibit an offstoichiometric compositional range.

15

Fig. 7 Calculated liquidus projection by Da Silva et al. [18]

To examine the solidification paths calculated by Da Silva et al. [18] different Alkemade lines are added in Figure 7. The first Alkemade line A1 is drawn between the VSS and the V5SiB2 phase and intersects with the respective (quasi-) binary eutectic valley between these two phases. Thus, the temperature decreases left and right along the two-phase eutectic valley as emphasized by the arrows which indicates a local maximum M′ (saddle point) along the VSS-V5SiB2 eutectic line. Consequently and in agreement with the Alkemade theorem, Da Silva et al. [18] identified two ternary eutectic reactions:

L ↔ VSS + V3B2 + V5SiB2

(E)

L ↔ VSS + V5SiB2 + V3Si

(E′)

The second eutectic automatically results also by drawing the Alkemade line A2 between the V3Si and V5SiB2 phase. Thus, a local maximum M″ along the V3Si-V5SiB2 eutectic valley is indicated. As a combination of the Alkemade lines A3 and A4 two additional maxima along the VB-V5SiB2 (M‴) and VB-V5Si3 (M‴′) mono-variant reactions

16

could exist. Hence, the calculated solidification path by Da Silva et al. [18] would emphasize a third ternary eutectic invariant reaction E″, respectively.

L ↔ V3Si + V5SiB2 + V5Si3

(E″)

The present experimental approach of the V-rich V-Si-B liquidus surface, however, would assume a slightly different solidification sequence as compared to Da Silva et al.’s [18] CALPHAD calculations. The experimentally determined liquidus projection is drawn with respect to the present model alloys, listed in Table 2, and their identified primary solidification areas as shown in Figure 8. In general, the present experimentally obtained results agree reasonably well with the liquidus surface predicted by the recent CALPHAD calculation [18]. Nevertheless, the main differences are found along the VSSV3Si, VSS-V3B2, VSS-V5SiB2, V3B2-V5SiB2 and V3B2-VB two-phase reaction lines and will be further discussed in more detail.

Fig. 8 Comparison of the V-Si-B liquidus projection calculated by Da Silva et al. [18] with the experimental results obtained in the present study. The primary phases of the

17

alloys investigated are indicated by symbols and compared to results on as-cast alloys taken from Reis et al. [17] (closed symbols) and De Lima [30] (open symbols).

As compared to the CALPHAD calculation [18] the eutectic VSS-V3B2 mono-variant line is found to be much longer than assumed by Da Silva et al. As a result, the twophase reaction line VSS-V5SiB2 is shifted to the right, towards higher Si-concentrations. Redrawing the Alkemade line A1 under these circumstances now reveals a different solidification sequence. The mono-variant VSS-V5SiB2 line did not intersect with A1 anymore and is located to its right hand side. Hence, a saddle point located at the VSSV5SiB2 eutectic line got impossible and the V3B2 phase decomposes in a class II (Utype) reaction into VSS and V5SiB2, while the ternary eutectic VSS-V5SiB2-V3Si (E′) could indeed be corroborated by the present experiments. Thus, the following U-type reaction is identified experimentally, which is in contrast to the calculations by Da Silva et al. [18] and their ternary eutectic reaction which was indicated as E in the present study:

L + V3B2 ↔ VSS + V5SiB2

(U)

The location of the V3B2-VB mono-variant line agrees well with Da Silva et al.’s [18] calculations

but

seems

to

be

slightly

extended.

The

U-type

reaction

L + VB ↔ V3B2 + V5SiB2 can be confirmed but seems to be slightly leaner in B and enriched in its Si-concentration. Thus, the V3B2-V5SiB2 mono-invariant line has a weak peritectic character and is considerably located towards higher Si-concentrations as compared to Da Silva et al. The size of the V3B2 primary solidification field seems to be almost double the size of the recent CALPHAD calculation [18]. A second difference of the present experimental study is addressed by taking a closer look on the VSS-V3Si binary eutectic reaction and its mono-variant line within the ternary system. The first approach on the V-Si-B liquidus surface [18] is using Zhang et al.’s [19] thermodynamic model of the V-Si binary boundary system. This model specially addresses the Si-rich part by key experiments on the Si-VSi2 eutectic reaction and the thermodynamic stability of V6Si5 [38]. While the calculated results fit reasonably well with experimental data for higher Si-concentrations (their Tab. 4 [19]), the discrepancy between their thermodynamic model and experimental results got bigger on the V-rich 18

portion of the V-Si binary system. Hereby, the VSS-V3Si binary eutectic plays a major role for the present study. Recent experimental results [25,26] revealed a binary eutectic point which has a slightly lower Si-concentration (12.6 at.%) than expressed by Smith’s binary phase diagram [23]. Da Silva et al.’s [18] recent calculation on the V-Si-B liquidus projection is based on Zhang et al.’s [19] thermodynamic model which emphasizes the eutectic point at 10.8 at% Si. The present experimental results, however, support the work by De Lima et al. [25] and Bei et al. [26] and is in agreement with high Siconcentrations for the VSS-V3Si eutectic reaction. Thus, its mono-variant line is also shifted towards higher Si-values as compared to the recent calculation [18]. Despite the aforementioned major differences of Da Silvas et al.’s approach and the present experimental investigations the results fit reasonably well at higher Siconcentrations. If, for example, the Alkemade line A2 between the stoichiometric V3Si and V5SiB2 phases is redrawn in Figure 8, the second ternary eutectic reaction L ↔ V3Si + V5SiB2 + V5Si3 (E″) can also be found. However, the present experimental study would assume a slightly lower B-concentration. Thus, the positions of the monovariant lines of V3Si-V5SiB2, V5SiB2-V5Si3 and VB-V5Si3 fit reasonably well with the CALPHAD calculation [18]. As a result of the present experimental investigations on the liquidus projection of the V-rich V-Si-B system, two class I ternary eutectic reactions (E-type) and three class II reactions (U-type) could be identified by the obtained microstructure observations of various cast V-Si-B alloys and by employing the Alkemade theorem. Figure 9 summarizes the solidification sequence by showing a partial reaction scheme of the Vrich part of the ternary system. It has to be mentioned here that the reactions are not indexed in their proper reaction order, since the V-rich corner of the V-Si-B system has an intermediate solidification temperature range and is thus, not the starting point of the reaction scheme in this system (which would be the B-rich corner due to the highest melting point of B in this system). However, similar to the Mo-Si-B system, the V-rich corner might be of major interest for possible future high temperature applications of VSi-B alloys and the present study contributes to a better understanding of the solidification behavior in this ternary system. Furthermore, the reaction temperatures are not known at present and the recent work by Da Silva et al. [18] can only give a rough estimation since their calculated solidification sequence might not corroborate with the 19

obtained experimental results. Extensive differential scanning calorimetry (DSC) work would be required to measure certain reaction temperatures which is quite challenging and cost intensive at temperatures higher than ≈ 1500 °C and might require special DSC equipment.

Fig. 9 Reaction scheme of the V-rich corner of the V-Si-B system obtained by the present experimental results. Information of the binary V-B and V-Si system are taken from the respective binary phase diagrams [23,24].

20

5. Summary and Conclusion

The present work is focused on an experimental approach to determine the V-Si-B liquidus projection in the V-rich corner of the ternary system. Hereby, recent CALPHAD calculations [18] were critical discussed using the Alkemade theorem and compared to the present experimental investigation and the microstructure evolution of various ternary V-Si-B alloys during cooling from the liquid state. The work can be summarized as follows: 1. The primary solidification fields and solidification paths in the V-rich V-Si-B system experimentally observed in this study are in good agreement with recent thermodynamic calculations by Da Silva et al. [18]. However, the size of the V3B2 primary solidification field differs strongly as compared to the calculations which further influences following cooling sequences. 2. Thus, the previous thermodynamic calculations [18] on the liquidus projection of the

V-Si-B

system

L ↔ VSS + V3B2 + V5SiB2

suggested (E),

three

ternary

eutectic

L ↔ VSS + V5SiB2 + V3Si

reactions: (E′)

and

L ↔ V3Si + V5SiB2 + V5Si3 (E″). The present experimental investigations can confirm the existence of only two ternary eutectic reactions, namely E′ and E″. 3. The

L ↔ VSS + V3B2 + V5SiB2

(E)

eutectic

reaction

could

not

be

found

experimentally. Thus, the present experiments indicate a weak peritectic transformation

L + V3B2 ↔ VSS + V5SiB2

(U-type)

preceding

the

L ↔ VSS + V5SiB2 + V3Si (E′) ternary eutectic. This conclusion is also in agreement with the Alkemade theorem and would then predict no saddle point along the VSSV5SiB2 eutectic valley (in contrast to the calculations by Da Silva et al. [18]). 4. The formation of the VSS-V5SiB2-V3Si ternary eutectic can now be described by the alloy’s solidification sequence. This will lead to further alloys development, since the Mo-Si-B neighbor systems share the similar type of ternary eutectic [31]. However, the VSS forms the major phases in the V-Si-B system which is favored in terms of low temperature mechanical properties [39].

21

Acknowledgements

The present research was funded by the German Research Foundation (DFG) under the grant number 410338871. Financial support of the Methodisch-Diagnostisches Zentrum

Werkstoffprüfung

(MDZWP)

e.V.,

Magdeburg,

Germany

is

greatly

acknowledged. The author kindly thanks M. Feuerbacher and C. Thomas (PGI-5, FZ Jülich) for providing access to the arc-melter, E. Wessel (IEL-2, FZ Jülich) for supporting the SEM analysis and M. Krüger (OVGU, Magdeburg), F. Stein (MPIE, Düsseldorf) and S. Kumar (Brown University, USA) for critical and fruitful discussions concerning the solidification paths of the V-Si-B system.

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25

Highlights -

The liquidus projection of the V-rich portion of the V-Si-B system has been experimentally investigated.

-

The solidification paths of alloys with primary VSS, V3B2, V3Si, V5SiB2, VB and V5Si3 phases were investigated.

-

Based on the experimental evaluation, a partial solidification scheme of the Vrich corner of the V-Si-B system could be achieved which differs from recant thermodynamic calculations.

-

Two ternary eutectic reactions, L ↔ VSS + V3B2 + V5SiB2 (Eʹ) and L ↔ V3Si + V5SiB2 + V5Si3 (E″), were observed and identified in the microstructures of V-rich V-Si-B alloys.

1 Author declaration

1. Conflict of Interest No conflict of interest exists. I confirm that there are no known conflicts of interest associated with this publication and there has been no significant financial support for this work that could have influenced its outcome.

2. Funding Funding was received for this work. All of the sources of funding for the work described in this publication are acknowledged below: German Research Foundation (DFG) under the grant number 410338871

3. Intellectual Property I confirm that I have given due consideration to the protection of intellectual property associated with this work and that there are no impediments to publication, including the timing of publication, with respect to intellectual property. In so doing I confirm that I have followed the regulations of our institutions concerning intellectual property.

4. Authorship I confirm that the manuscript has been read and approved by all named authors. I confirm that the order of authors listed in the manuscript has been approved by all named authors. I the undersigned agree with all of the above.

Author’s name (Fist, Last) 1. Georg Hasemann

Signature

Date 10.25.2019