The Detroit Post
Thursday, 25 February, 2021

Are There Twins In Bitlife

David Lawrence
• Tuesday, 03 November, 2020
• 14 min read

You have a lot of freedom in Billie, the mobile game simulator, where you can do almost anything with zero consequences. You can even find yourself going to jail, where you later have to break out or modifying the settings in God Mode, where you can change a person’s willpower.

triplets comments bitlifeapp


For those who are looking to have twins, you can roll the dice by naturally going through the process of having a child with your partner. The option for IVF shows up if you and your partner are struggling to have a child in the game.

If they do agree by going several conversations, you and your partner have a much higher chance of your offspring being twins or even triplets. If you’re attempting to have any twins or triplets, though, continue down these two routes to make it possible.

In this tutorial, we will show you how to have twins and triplets in the Billie game. Once you have graduated from college and have found a perfect job that might meet your desired necessities, then you might consider settling down in your life.

But rather than just one single addition, some people are looking for ways to make welcome two or three new members to their family. But the ability to have twins and triplets in Billie is proving to be an altogether different challenge in itself.

Today, we will be listing out all the requirements that need to be met from you and your partner’s end to achieve the aforementioned quest. The first one is having the birth the normal way and hoping that it might be a twin or a triplet.

However, that is completely based on luck with the probability on the lower side. Some users have even tried this for continuously three-four eyes but weren’t able to achieve success.

While both these methods have a much higher chance of offering twins and triplets in Billie, it comes with its fair share of issues as well. If it fails to give out the desirable results, then you will have to wait for a year and then retry it.

As we have already covered several aspects of the game before, we will entirely be focusing on having siblings now and how having brothers and sisters can impact your virtual life. It’s very typical as well to have only 1 brother or sister but if you play the game even for just a short while, you would be surprised as to how fast your family can grow.

Though it would still depend on probabilities, there ’s a big chance to have brothers or sisters close to your age so if you will be having younger siblings, they should start popping out between 2 and 4 in-game years. Unless you are a real toddler yourself or practically asking for negative repercussions and bad karma, you can pick either one and reduce both your and your sibling’s level of happiness.

That won’t be the worst of it though, as any of these choices can lead to violent retaliations, injuries, or possibly even punishments from your parents. As a third, and obviously positive option, you can choose to do nothing and wait for your siblings to grow up a bit more so you can than begin having conversations or spending some time with them.

You can always try to assault or engage in petty fights with them but it will quickly lead to you ending up with an empty level of happiness that would be hard to get out of. On top of all of this, doing bad things to your siblings may cause them to hurt you, and while there doesn’t seem to be any sort of permanent damage, it can be very difficult to recover from.

The events may not be entirely related to the person you hurt but some things tend to go wrong randomly when you are in bad terms with a lot of people whom you should be keeping a good relationship with. Engaging in an argument or petty squabble with a sibling with a full petulance level over and over won’t risk you being hurt back by them.

One of the biggest advantages of having more siblings is that in cases of unfortunate events that leave you unhappy, you can easily gain back a lot of happiness just by spending time with each of them. While keeping a good relationship with everyone in a huge family may take more time and effort, it can be especially rewarding as you will hardly worry about depression and anxiety even in the worst series of hardships that you experience in your virtual life.

Engaging in only positive activities with your siblings will lead to a much simpler life as they will serve to be a constant happiness booster for you themselves and even more once they bear their own children. Simply telling your parents about it may lessen their level of happiness on top of a variety of possible actions they may do as an effect of their children’s squabbles.

If you want a simple life with siblings that dwell on constant happiness and good karma, then stick with what you feel would be the positive actions and reactions. As there are other new contents brought about by the recent update, there may still be some more unique elements that would affect your virtual life relative to having siblings.

We hope you were able to read through our short and simple guide and that you have learned important tips and strategies you can use as you play the game. 2 Mechanical and Materials Engineering, University of Nebraska-Lincoln, Lincoln, NE 68583, USA.

Twinning is a critically important deformation mode in hexagonal close-packed metals. An understanding of twin transformations therefore necessitates that the atomic-scale structure and intrinsic mobilities of facets be known and characterized.

The present work addresses the former point by systematically characterizing the boundary structures of 3D {1¯012} twins in magnesium using high-resolution transmission electron microscopy (HR TEM). Eight characteristic facets associated with twin boundaries are reported, five of which have never been experimentally observed before.

Further, molecular dynamics simulations suggest that the formation and motion of these facets is associated with the accumulation of twinning dislocations. Twinning is a major plastic deformation mechanism accommodating shear in hexagonal close packed (HCP) and other lower symmetry metals (i.e., Mg, Ti, Zr, Be, Sn, and U) (1 – 3), as well as in some high-performance steels (4, 5) and in many cubic metals during high rate/stress loading (6, 7).

This mechanism has been discussed in a series of recent animistic simulation studies focused on the migration of two such facets: the coherent twin boundary (CTB) and basal-prismatic (BP/PB) (11 – 16). It has been recently shown that TD's and steps/facets are favorable sites for segregating impurities and alloying elements (14, 17 – 19).

Specifically, while twin domains are necessarily bound by numerous facets each with distinct geometry and structure, previous experimental studies have defined the twin boundary structure only from a single direction and thus elucidate only a single “boundary character.” This would be akin to having an understanding of dislocation-mediated plasticity limited to considerations of only one dislocation character. At the micrometer scale, however, twin boundaries generally deviate significantly from the CTB relationship (16, 24 – 26).

Experimentally, the atomic-scale defects defining twin boundaries have only been definitively characterized along a single crystallographic direction, the <12¯10>, , that lies in the K 1 plane and is perpendicular to 1 (8, 11, 16, 25 – 29). BP/PB facets closely align (0001) and (101¯0) planes in the twin and parent, respectively, and K 2 facets closely align the {1¯012¯}{1¯012¯} K 2 planes in the twin and parent.

There is only one previous experimental study at the atomic scale of a {1¯012} twin boundary viewed along a different direction, the <101¯1>, 1 (30). The study relied on complementary molecular dynamics (MD) simulations to suggest the serrated boundary to be composed of Cabs and semi coherent twist prismatic-prismatic (Twist-Pr2Pr2) facets with {2¯110}{211¯0}.

Considering that twin growth takes place in 3D, there is a fundamental need for elucidating the propagation of twin interfaces in 3D, to which the studies from the <12¯10> direction provide only a partial answer. Here, we tackle the first and essential step in such a quest, namely, figuring out what are the facets that bound twin domains and their atomic configurations.

We have undertaken the experimental atomic-scale characterization of the most common twin in HCP metals from six crystallographic directions and have found evidence of eight different boundary facets, five of which have never been experimentally observed. MD simulations further suggest that the formation and motion of these facets are associated with the accumulation of TD's.

Atomic-scale imaging of facets in TEM foils requires specific imaging conditions related to the foil normal, facet geometry, and crystallography. To image identifiable facets at the atomic scale, three conditions must be simultaneously satisfied.

The crystallographic observation directions depend on the atomic configuration of the twinning plane, which is shown in Fig. While the four-index Miller-Bravais indices for the observation directions seem high, the three-index Miller index equivalents are low, with <54¯1¯3> and <422¯3> being <21¯1> and <201>, respectively.

Second, because of the tilt limitation of the TEM holder, the deviation angle between the imaging direction and the normal direction of the TEM foil must be within the tilt range (<30°). We performed electron back scattered diffraction (BSD) analysis on twin-jet polished TEM foils to preselect twins with proper orientations.

In many cases, twin facets in a TEM foil are inclined relative to the imaging direction resulting in a twin boundary with a diffuse appearance, and it is not possible to define the facets from the specific imaging condition. There is an experimental analogy between HR TEM imaging of twin facets and dislocation cores.

We report only the twin facets that are identifiable for the given TEM foil and imaging conditions. (B) Schematic of the facets identified from {1¯012} twins observed by HR TEM from the six low-index crystallographic directions indicated in (A) and (B).

A 3D twin domain of arbitrary shape could be built by combining the basic elements of Cabs and these facets in different length/height ratios. Attempts were made to perform HR TEM at higher index zone axes, but none of the higher index zone axes showed facets and, in addition, gave relatively poor HR TEM imaging conditions.

A 3D twin domain with an arbitrary shape could be formed by combining the basic elements of Cabs and these facets in different length/height ratios. For instance, the well-defined BP interface has a 3.7° angle between the basal planes from the matrix/twin and the prismatic planes from the twin/matrix for {101¯2} twins in Mg. We use simulated overlapping diffraction patterns for each zone axis to help visualize the planes in the matrix and twin that are nearly parallel and likely candidates to form interfaces.

Some diffraction spots from the matrix and twin nearly overlap, indicating that the planes are nearly parallel. Second, we combined different methods of identifying the edge-on facets in the HR TEM images depending on the observation zone axis.

Diffraction contrast is a conventional TEM contrast mechanism often used at magnifications less than around 100,000× relying on the fact that the matrix and twin have different crystal orientations and, thus, diffract electrons differently for most imaging conditions. GPA is normally used to measure in-plane displacements and/or strains at the nanometer scale within an HR TEM field of view using a fast Fourier transform (FFT)–based analysis of local lattice parameters (31, 32).

We use GPA in this study to measure small magnitude plane reorientation at the twin boundary for the <54¯1¯3> zone axis. The <101¯1> zone axis represents a special case where the mirror image of the families of planes in the matrix exactly aligns with the same planes in the matrix, and thus, the boundary position cannot be determined based on a mirror symmetry relationship.

However, the atoms within the interface are at slightly different positions relative to the parent and twin lattices, and the interface plane is often evident due to the resulting HR TEM phase contrast. Details of the methods used for determining boundary positions for each zone axis are given in the following content.

Figure 2D shows the atomic structure of a BP facet with dislocations labeled in the boundary. On the basis of the atomic structure, we studied the formation and stability of long BP facets in Mg in a recent paper (33).

2Simulated diffraction patterns and images from {1¯012} twins viewed along a <12¯10> zone axis within a K 1 plane. (B) HR TEM image from an extended twin boundary containing Cabs and a BP facet.

(C) HR TEM image from twin tip containing Cabs and BP facets. The inset is the selected-area electron diffraction (Said) pattern from both the matrix and twin.

For the <541¯3> zone axis, the simulated overlapping diffraction patterns shown in Fig. 3Simulated diffraction patterns and images from {1¯012} twins viewed along a <54¯1¯3> zone axis within a K 1 plane.

(C) GPA result for the square area in (B) overlapped with the corresponding HR TEM image, showing the relative rotation of {01¯11¯} planes from twin to matrix. (E) HR TEM image from an extended twin boundary containing Cabs and PyPy1 facet.

(F) GPA result for the image in (E), showing the relative rotation of {01¯11¯} planes from twin to matrix. Figure 3B shows the overall shape of a twin tip in an HR TEM micrograph viewed along a <54¯1¯3> direction in the K 1 plane.

3B shows the selected-area electron diffraction (Said) pattern from both the matrix and twin. 3B, FFT analysis was performed on the tip area (shown in fig.

The areas with only one set of diffraction spots (squares 1 and 3 in fig. The colored map shows the relative rotation of {01¯11¯} planes from twin to matrix.

MD simulations in what follows predict the existence of the PyPy1 facets and support this experimental result. An HR TEM image from an extended twin boundary is shown in Fig.

It shows a step along the CTB, which is also shown in the inset without lines for clearer observation purposes. The simulated overlapping diffraction patterns for a <422¯3> zone axis shown in Fig.

4Simulated diffraction patterns and images from {1¯012} twins viewed along a <422¯3> zone axis within a K 1 plane. © Enlarged HR TEM image of the square area in (B) showing the {01¯10}{1¯21¯2} and {14¯32¯}{1¯43¯2} facets.

(D) HR TEM image from an extended twin boundary containing CTB and {01¯10}{1¯21¯2} facets. The interfaces can be identified as {01¯10}{1¯21¯2} and {14¯32¯}{1¯43¯2}, which are indicated by yellow and red dashed lines (the identification details are shown in fig.

The diffraction spots from the matrix (black dots) and twin (red dots) in the <101¯1> shear direction exactly overlap, as shown in the simulated overlapping diffraction patterns of Fig. Because the diffraction patterns of the twin and matrix exactly overlap, there are many parallel planes, the first three of which (as ordered by d-spacing) are (1¯101) M (01¯11) T, (0111¯) M (11¯01¯) T, and (1¯21¯0) M (12¯10) T.

5Simulated diffraction patterns and images from a {1¯012} twin viewed along the <101¯1> shear direction. (A) Simulated overlapping diffraction patterns; all diffraction spots from the twin (red dots) overlap with the spots from the matrix (black dots).

(D) Enlarged image of the square area in (C), showing the Cabs and PyPy1 facet at the step. Figure 5B shows a low-magnification image of a {1¯012} twin viewed slightly off of the <101¯1> shear direction.

5C is obtained by zooming in on the extended twin boundary indicated in Fig. While this is not helpful for distinguishing the atomic position of the boundary, it does confirm that the twin boundary defects viewed along the shear direction do not have significant edge character.

The bright spots within the boundary are due to HR TEM phase contrast. The groups of {1¯101} planes in matrix and twin are labeled by yellow and red dashed lines in Fig.

The boundaries inside the rectangle exhibit contrast and are enlarged in Fig. According to the simulated overlapping diffraction patterns, it is likely a prismatic-(1¯21¯0) M prismatic-(12¯10) T (PrPr2) interface, which is perpendicular to the CTB.

We summarize the facets defining {1¯012} twin boundaries determined by HR TEM in Fig. In what follows, we use MD simulations for two purposes: (i) to substantiate the 3D animistic configuration of the facets revealed by the essentially 2D sectioning process and (ii) to lay the groundwork for future animistic studies of the mobility of these facets.

Unless agreement between measured and predicted equilibrium facets is achieved, one cannot take the next step consisting in dynamic propagation simulations. Figure 6A shows the initial twin structure at 100 K. A 1-GPa resolved shear stress is imposed on the twin system to grow the twin domain while maintaining the temperature of the overall domain at 100 K. Details of the MD construction and simulation are described in Materials and Methods and in the Supplementary Materials.

In the simulation, Cabs become the largest interfaces through propagation of existing TD's. When viewed along the direction, BP/PB and K 2 facets (delineated by solid red lines in Fig.

6B) are the second largest facets observed, which is consistent with the TEM images in Fig. Viewed along the direction, {011¯1}{01¯11¯} facets (delineated by red solid lines in Fig.

The corresponding facets observed in the experiments are highlighted by yellow dashed lines in Fig. When viewed along the direction, small {01¯10}{1¯21¯2} facets (colored green in Fig.

6A) are observed, consistent with the cross-section view normal to the direction in Fig. 6D, a small vertical segment along the trace of the {14¯32¯}{1¯43¯2} interface on the cross-section is circled.

However, this short segment is not the small facet depicted by green dashed line in Fig. 4A but the projection of {011¯1}{01¯11¯} facets along the direction, as can be inferred when comparing to the 3D configuration in Fig.

3A suggests a relatively high formation energy of this interface, and the MD simulation may not capture the feature. As a consequence, the twin domain not only must grow in thickness and propagate forward in the shear direction but also needs to expand laterally.

The growth process is controlled by the formation energy and mobility of the 3D twin facets discovered here. MD allows one to address the kinetics evolution during twin growth, and our results indicate that this process is strongly anisotropic.

The overall agreement between the predicted and the experimentally observed twin facets suggests that MD provides a viable path to predict and rationalize observed 3D microstructures. For instance, the simulations here indicate that twin propagation is not symmetric: Twins propagate faster in the lateral direction relative to the “forward” shear direction.

The overall shape of the twin, wider along the lateral dimension than along the propagation direction, is consistent with experimental BSD observations reported by Liu et al. These cubical samples were compressed to a total engineering strain of 1.3% along the 10-mm direction, approximately perpendicular to the basal pole for most grains, resulting in twinning activity in most grains.

For the first orientation, the foil normal is perpendicular to the compression and TT directions such that zone axes within 45° of the <112¯0> are accessible for many twins. For the second orientation, the foil normal is around 45° to the compression and TT directions such that zone axes within 45° of the 1 <101¯1> are accessible for many twins.

The 3-mm discs were electropolished to perforation in a solution of 2% nitric acid and water, a voltage of 0.1 V, and temperature of 2 °C. Twin orientations and tips of interest were preselected using BSD in an FEI Are scanning electron microscope.

The simulations use the modified embedded-atom method (Team) potential developed for magnesium by Wu et al. Figure S5A shows the coherent chromatic pattern associated with pure-shuffle nucleation of a {1¯012} twin.

S5D is obtained at 5 K with multiple loading-unloading cycles under a deformation gradient with one nonzero simple shear component F 12. The twin mainly propagates along the direction and develops into a structure consisting of multiple facets.

Acknowledgments: We thank K. Dang for creating the supplementary animistic simulation sections of a {1¯012} twin boundary showing the slightly shifted atoms at twin boundaries. Funding: This work is fully supported by the Office of Basic Energy Sciences, Project FDP 06SCPE401, under U.S. DOE contract no.

Additional data related to this paper may be requested from the authors. Copyright © 2020 The Authors, some rights reserved; exclusive licensee American Association for the Advancement of Science.

Other Articles You Might Be Interested In

01: Alabama Real Estate Montgomery Al
02: Alamo Ranch Restaurants San Antonio Tx
03: Alaska Academy Of Real Estate Anchorage Ak
04: Albuquerque Commercial Real Estate For Lease
05: Albuquerque Craigslist Real Estate For Sale
06: Albuquerque Journal Best Restaurants
07: Albuquerque Journal Real Estate
08: Albuquerque Mls Real Estate Listings
09: Albuquerque Real Estate Brokers
10: Albuquerque Real Estate By Owner Craigslist
1 -
2 -
3 -
4 -
5 -
6 -
7 -
8 -
9 -